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
123
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
[email protected] (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.
[email protected] 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.
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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 <110> 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
[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 (<10 nm) and which has a small convergence angle so that the diffraction patterns show sharp spots. To obtain this condition, a small condenser aperture (10 μm) is used in microprobe STEM mode in a Tecnai F30. Two series of transistors, having Si80Ge20 and Si99C1 S/D respectively and different channel lengths were investigated. These epitaxial layers were grown using an ASM Epsilon® 2000 reactor. TEM specimens were prepared by focused ion beam for [1-10] zone axis observation. Series of electron diffraction patterns were acquired in the channel region perpendicular to the gate (Figure 1). Data were processed using the FEI True Crystal Strain software which automatically determines the center of the diffraction spots and calculates the displacement of their positions with respect to the reference diffraction pattern. Strain variations in the horizontal [110] and perpendicular [001] (growth) direction, were obtained along a line profile perpendicular to the interface. Experimental data were converted from strain to stress and then compared with simulations performed by the finite-element simulator Taurus-Process by Synopsys [3]. Figure 2 shows the horizontal strain variation measured perpendicular to the interface in channels having on mask lengths 500nm, 250nm and 90nm respectively and 250nm wide Si1-xCx S/D. As expected Si1-xCx induces channel tension. The strain is smaller (0-0.2%) in larger channels and reaches higher values of about 0.5% in the 90nm channel. The strained region extends deeper in the Si substrate as the channel length is reduced.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 15–16, DOI: 10.1007/978-3-540-85226-1_8, © Springer-Verlag Berlin Heidelberg 2008
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Figure 3 shows the horizontal experimental and simulated stress variation perpendicular to the interface in structures having Si channels and Si1-xGex S/D. In these devices the stress is, as expected, compressive in the channel along the [110] direction. Measured stress is in good agreement with the simulated one. Table 1 shows a summary of the experimental results for maximum strain and depth of strain penetration in the Si substrate as a function of the gate length for the two types of investigated devices. In general, the shorter the gate, the larger is the strain (either compressive or tensile) and the deeper is the penetration in the Si substrate. The results obtained in this work by NBD are fairly consistent with theory and simulations, proving that the NBD technique can be reliably applied to such structures. 1. 2. 3.
B. Foran, M. H. Clark and G. Lian, Future Fab International 20 (2006), p. 127. N. Bich Yen, V. Vartanian, A. Thean, et al., Solid State Technology 49 (2006), p. 41. Taurus TSUPREM, Taurus Process Reference Manual, Oct. 2005. Version X-2005.10. a
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Figure 3. Experimental (marker) and simulated (line) stress along [110] in Si channels with SiGe S/D. The lineprofile is perpendicular to the interface (Figure 1). Table I S/D SiC SiC SiC SiGe SiGe SiGe
On mask gate length (nm) 500 250 90 200 130 100
Max perpendicular strain (%) 0.1 0.3 0.5 -0.8 -1.3 -1.8
Strain penetration in the Si (nm) 40 90 130 150 100 100
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Cluster growth and luminescence in ion-implanted silica H.-J. Fitting1, Roushdey Salh1, L. Kourkoutis2, and B. Schmidt3 1. Physics Dept., Rostock University, Universitaetsplatz 3, D-18051 Rostock, Germany 2. Applied and Engineering Physics, Cornell University, Ithaca NY 14853, U.S.A. 3. Institute of Ion Beam Physics, Res. Center Dresden-Rossendorf, D-01314 Dresden, Germany
[email protected] Keywords: Cathodoluminescence, Silica layers, Ion implantation
Scanning electron microscopy (SEM) and cathodoluminescence (CL) in combination with scanning transmission electron microscopy (STEM) have been used to investigate thermally grown amorphous silicon dioxide layers implanted isoelectronically with group IV ions (C+, Si+, Ge+, Sn+, Pb+) as well as with group VI ions (O+, S+, Se+), see Fig.1. As the main experimental results we can state: Oxygen surplus increases the red band (R) in SiO2 but does not affect the blue band (B). Silicon surplus increases the blue luminescence B, but reduces the red band (R) as already demonstrated in [1]. So it is verified that the red luminescence (R) is an oxygen related center stated as NBOHC and the blue luminescence (B) is a silicon related oxygen deficient center Si-ODC. In Ge+ implanted SiO2 a strong violet band V (3.1 eV ; 410 nm) appears and is associated with the Ge-related oxygen deficient center Ge-ODC. This V band shows a huge increase after thermal annealing of Ta=900 °C by more than two orders of magnitude. Thus the Si-ODC's as well as the Ge-ODC's are formed by Si and Ge molecules clustering to certain low-dimension fragments like dimers ODC-I, trimers and up to higher aggregates, [2]. The nanocluster size grows with annealing temperature from 2-4 nm at Ta=900 °C to 5-10 nm at Ta=1100 °C, see e.g. Fig. 2. Further ion implantations of group IV elements: C, Ge, Sn, Pb which are thought to substitute Si isoelectronically in the silica matrix, show new bands and a general increase of the luminescence in the violet-blue region. On the other hand, implantations of oxygen substituting elements of group VI (S, Se) lead to an increase in the yellow-red spectral region, [2]. As a surprising peculiarity, the cathodoluminescence spectra of the group VI elements oxygen and sulfur implanted SiO2 layers show, besides characteristic bands, a sharp and intensive multimodal structure as shown in Fig. 3. The energy step differences of the sublevels amount to an average of 120 meV and indicate vibronicelectronic transitions, probably, of O2− interstitial molecules, as we could demonstrate by a respective configuration coordinate model, [3]. 1. 2.
H.-J. Fitting, T. Barfels, A. N. Trukhin, B. Schmidt, A. Gulans, and A. von Czarnowski, J. Non-Cryst. Solids 303 (2002) 218 – 231. Roushdey Salh, L. Fitting-Kourkoutis, B. Schmidt, H.-J. Fitting, Physica Status Solidi (a) 204 (2007) 3132 – 3144
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 17–18, DOI: 10.1007/978-3-540-85226-1_9, © Springer-Verlag Berlin Heidelberg 2008
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H.-J. Fitting, Roushdey Salh, T. Barfels, and B. Schmidt, Physica Status Solidi (a) 202 (2005) R142 – R144. 80
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Figure 1. Electron beam excitation in SiO2 layers for different beam energies Eo allowing a CL depth profiling (left). Here a Ge profile is implanted in the mean projected range Rp ≈250 nm. The STEM micrograph (contrast inverted) shows the actual Ge cluster profile.
Figure 2. Ge cluster growth in silica with increasing annealing temperatures Ta, (Ostwald ripening).
Figure 3. Multimodal CL spectra from sulfur-implanted SiO2:S layers, recorded at liquid nitrogen temperature (LNT).
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Coherence Measurements of Bulk and Surface Plasmons in Semiconductors by Diffracted Beam Holography R.A. Herring Center for Advanced Materials & Related Technologies, Mechanical Engineering, University of Victoria, Victoria Canada V8W 3P6
[email protected] Keywords: plasmon loss electrons, electron holography, coherence
The degree of coherence (γ) and lateral coherence (δ) of plasmon loss electrons has been measured using two methods of energy-filtered Diffracted Beam Interferometry/Holography (DBI/H); one using high-angle scattered electrons produced by a small probe of a few Angstroms referred to as convergent DBI/H [1] and another one using a variable, large-diameter planar beam referred to as planar DBI/H [2] (Figure 1). Both DBI/H methods require a (S)TEM with an electron source having good coherence, an electron biprism for holography and an imaging energy filter (GIF) for separating out the plasmon loss electrons. Surface and bulk plasmon excitations are created by the fast electron during its passage into and through the material specimen. After the creation event, these plasmon loss electrons retain all of the plasmon’s coherence information. An important consideration in these measurements is dynamic diffraction, which is accounted for in the simulations of energy-filtered convergent DBI/H by use of the mixed dynamic form factor [5]. To help reduce dynamic diffraction, thin specimens of ~¼ extinction distance of the diffracted beam (~15 nm) have been used, which also produces equal intensity beams that provide for the highest contrast fringes and simplifies the measurement of γ and δ. Additionally, convergent DBI/H uses the condenser aperture focused on the diffraction plane for separating out the source electrons from the plasmon loss electrons at large scattering angles (Figure 2). A γ of ~0.1 was measured at scattering angle of ~ 7 mrad. The fringes disappeared when the interfering beams were separated using the electron biprism from the exact overlay position, which has produced measurements of the lateral coherence of ~1 nm on the specimen plane. It is believed that the surface plasmon loss electrons dominate the energy-filtered electron intensity at large beam diameters. As the planar beam diameter becomes larger in planar DBI/H, the surface plasmons lose their mutual coherence resulting in a decrease and eventual loss of their fringe contrast in the center of the interferogram giving a measurement of δ (Figure 3). Knowledge of δ of bulk plasmons and surface plasmons would be very helpful for understanding nanoscience and for developing nanotechnology. 1. 2. 3.
R.A. Herring, Ultramicroscopy 104 (2005) 261. R.A. Herring, Ultramicroscopy (2007). In press. P. Schattschneider and J. Verbeeck, Ultramicroscopy (2007). In press.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 19–20, DOI: 10.1007/978-3-540-85226-1_10, © Springer-Verlag Berlin Heidelberg 2008
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FEI, Eindhoven and CEMES, Toulouse are gratefully acknowledged for the use of their microscopes. Grants from UVic, NSERC and the CFI are also gratefully appreciated.
a) b) Figure 1. Schematics showing convergent DBI/H in a), an amplitude splitting method with no interference of the beams and biprism, and planar DBI/H in b), a mixed amplitude plus wavefront spitting method having beam and biprism interference.
a)
c) b) Figure 2. Energy-filtered convergent DBI/H of plasmon loss electrons of Ge showing an over-exposed hologram in a), intensity profile extending away from center along white line in b), and retained coherence of high-angle scattered electrons in c).
γ = 0.2
d) a) c) Figure 3. Energy-filtered convergent DBI/H of plasmon loss electrons of Ge (in a) and b)) and of Si (in c) showing a hologram at position 1 in Fig 2a using a small electron probe of a few Å in a), and planar DBI/H holograms in b) using 200 nm diameter beam and in c) using 2000 nm diameter beam, giving δ ~ 2000 nm for Si.
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Comparison of 3D potential structures at different pn-junctions in FIB-prepared silicon and germanium samples measured by electron-holographic tomography A. Lenk1, D. Wolf1, H. Lichte1 and U. Mühle2 1. Triebenberg Laboratory, Institute of Structure Physics, Technische Universität Dresden, D-01062 Dresden, Germany 2. Qimonda Dresden GmbH & Co. OHG, Königsbrücker Straße 180, D-01099 Dresden, Germany
[email protected] Keywords: p-n junction, intrinsic electrostatic potential, specimen surface
The functionality of electronic nanodevices is essentially determined by location, extension and concentration of the enclosed doped areas. Therefore, the manufacturing of those areas is one of the most important tasks in semiconductor industry. In general, it consists of two steps: First, the dopant material has to be inserted into the lowly doped semiconductor substrate. Then, the inserted dopants have to become electrically activated. Electron holography allows to monitor the outcome of both processes simultanously by providing 2D mapping of the resulting electrostatic potential distribution in the semiconductor [1]. Though, prerequisite for such an analysis is an appropriate preparation method. Precise target preparation, needed for the very small semiconductor devices, leaves no alternative to the application of a Focussed-Ion-Beam (FIB) tool. It has been shown that the method can be adapted to the special requirements of electron holography [2]. Unfortunately, the FIB-treatment always alters the electrical and structural characteristics of the sample, or at least those of the micromachined lateral surfaces and their adjacent regions. For a successful quantitative evaluation of the measured data, this has to be understood and taken into account. For this purpose, a new needle-shaped specimen geometry (shown in Figure 1) was designed and used to investigate a conventionally FIB-prepared lamella in cross-section [3]. Such needles are also highly qualified for tomographic application in the TEM. So it was possible to measure and compare 3D distributions of the electrostatic potential in silicon needle samples with pn-junctions of both types. Figure 2a/b shows a comparison of two virtual slices through the inner core of reconstructed 3D needles. It was found out, that the potential ascend in n-doped areas surrounded by p-doped substrate (0.44 V in Figure 2a) is smaller then the potential descend in p-doped regions surrounded by ndoped substrate (0.92 V in Figure 2b) [4]. In theory, both differences should be equal. To investigate this effect further, a germanium-based needle with an n-doped top layer has been analysed with electron holography, conventionally and tomographically. Figure 2c shows the holographically measured potential shift at the germanium pnjunction. Again, the measured potential drop ist lower than theoretically expected.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 21–22, DOI: 10.1007/978-3-540-85226-1_11, © Springer-Verlag Berlin Heidelberg 2008
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1. 2. 3. 4.
W.-D. Rau, P. Schwander, F. H. Baumann, W. Höppner, and A. Ourmazd, Physical Review Letter 82 (1999), p. 2614. A. Lenk, H. Lichte and U. Mühle, Journal of Electron Microscopy 54(4) (2005), p. 351. A. Lenk, U. Mühle and H. Lichte, Springer Conference Series “Microscopy of Semiconducting Materials XIV”(2005), p. 213. A. Lenk, D. Wolf and U. Mühle, Conference Proceedings of MC Saarbrücken (2007), p. 306.
Figure 1: Images of a needle-shaped specimen. Left: Top view in the FIB after cutting. Right: Bright field image in a TEM. Because of its special symmetric geometry, the needle can be used as a cross-section of a conventionally FIB-prepared lamella.
Figure 2. Holographically measured potential differences at different pn-junctions in silicon and germanium samples. 2a: n-doped layer in p-doped silicon. 2b: p-doped layer on n-doped silicon. 2c: n-doped layer on p-doped germanium. The images 2a & 2b are slices from tomographically reconstructed virtual needles, whereas image 2c is a phase image from a conventionally recorded hologram.
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EELS analyses of metal-inserted high-k dielectric stacks M. MacKenzie1, A.J. Craven1, D.W. McComb2, C.M. McGilvery2, S. McFadzean1 and S. De Gendt3 1. Department of Physics & Astronomy, University of Glasgow, Glasgow, G12 8QQ, UK 2. Department of Materials, Imperial College London, London, SW7 2AZ, UK 3. IMEC, Kapeldreef 75, B-3001 Leuven, Belgium. Also at KU Leuven, B-3001 Leuven, Belgium
[email protected] Keywords: high-k dielectric, metal gate, EELS, ELNES
Hf-based systems are going into production this year as the high-k materials replacing amorphous SiO2 and Si(O,N) as the gate dielectric in Si MOSFETs. At the same time, metal inserted poly-Si gate electrodes are being used in the gate stacks to remove problems associated with poly-Si gate electrodes. Significant interface interactions can occur in such systems as a result of the thermal budget received during device processing. We are investigating a range of HfO2 and HfSiO based stacks with TiN or TaN metal inserted poly gates. The effect of different deposition methods and processing treatments on physical and chemical properties are probed via electron energy loss spectroscopy (EELS). In particular, interface reactions occurring between the layers in the stack are investigated using the electron energy loss near edge structure (ELNES) to extract information on the phases present. The specimens are examined in an FEI Tecnai F20 TEM/STEM equipped with a field emission gun, a Gatan ENFINA electron spectrometer and an EDAX X-ray spectrometer. Spectrum imaging was performed using Gatan DigiScan II and Digital Micrograph software. Mounted on one of the 35 mm camera ports is a retractable fast beam switch (FBS) [1,2]. The FBS allows us to collect the EELS core and low loss spectra from the same regions under the same electron-optical conditions and hence to remove the effects of plural scattering and normalize the overall signal level. Absolute elemental quantification is thus enabled. Figure 1(a) shows a high resolution TEM image of a Si/SiO2/HfSiO/TaN/poly-Si gate stack. As well as an amorphous phase at the TaN/poly-Si interface, one can observe lighter patches in the TaN layer. The high angle annular dark field (HAADF) STEM image of the same stack in Figure 1(b) shows dark notches penetrating into the TaN from the TaN/poly-Si interface. We have previously reported some EELS results across this stack and shown the presence of an oxide phase at the TaN/poly-Si [3]. Here we investigate the dark notches in the HAADF STEM images. An EELS spectrum image was acquired across one of these notches along the white line in Figure 1(b). Figure 2(a) contains the N and O elemental profiles obtained from the spectrum image and shows that the dark notch corresponds to an oxidised region of TaN. The
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 23–24, DOI: 10.1007/978-3-540-85226-1_12, © Springer-Verlag Berlin Heidelberg 2008
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background subtracted O K-edge part of the spectrum image shown in Figure 2(b) also confirms this.
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Figure 1. (a) High-resolution TEM image of a Si/SiO2/HfSiO/TaN/poly-Si gate stack. (b) HAADF STEM image of a Si/SiO2/HfSiO/TaN/poly-Si gate stack. An EELS spectrum image dataset was acquired across the gate stack as indicated by the white line.
Energy Loss (eV)
Figure 2. (a) N and O elemental profiles extracted from an EELS spectrum image acquired along the line shown in Fig. 1(b); the substrate is on the LHS. (b) Background subtracted O K-edge spectrum image. The Si substrate is at the top. 1. 2. 3. 4.
A.J. Craven, J. Wilson and W.A.P. Nicholson, Ultramicroscopy 92 (2002), p. 165. J. Scott, P.J. Thomas, M. MacKenzie, S. McFadzean, J. Wilbrink, A.J. Craven and W.A.P. Nicholson, submitted to Ultramicroscopy (2008). M. MacKenzie, A.J. Craven, D.W. McComb, C.M. McGilvery, S. McFadzean and S. De Gendt, submitted to proceedings of MSMXV 2007, Cambridge. We kindly acknowledge the EPSRC for funding (GR/S44280) and B. Miller (University of Glasgow) for TEM specimen preparation.
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Low voltage SEM observations of the dopant contrast in semiconductors K. Masenelli-Varlot1, S. Luca1, G. Thollet1, P.H. Jouneau2 and D. Mariolle3 1. Université de Lyon, MATEIS UMR 5510, INSA-Lyon, 7 avenue Jean Capelle, 69621 Villeurbanne, France 2. CEA/INAC/SP2M/LEMMA, 17 avenue des Martyrs, 38054 Grenoble, France 3. CEA/LETI, 17 avenue des Martyrs, 38054 Grenoble, France
[email protected] Keywords: semiconductor, SEM, doping contrast
The electrical properties and performances of most semiconductor devices depend on their dopant concentration. For this reason the Semiconductor Industry needs a twoand three-dimensional dopant profiling technique with nanometric spatial resolution in the range of sensitivity between 1016 and 1020 atm/cm3 and an accuracy of ±10%. One of the few 2D dopant profiling technique that can fulfil these requirements is the scanning electron microscopy (SEM), which is a fast and non-destructive technique, with a rapid sample preparation. However it was shown that it has some limitations because of the lack of reproducibility in the image acquisition [1]. In this paper we present a study on the doping contrast in Si semiconductor structures by Low Voltage Scanning Electron Microscopy (LVSEM). The experiments were performed on low vacuum FEG-SEM Zeiss Ultra55 and FEI XL 30 ESEM. The particularity of the last one is that it can work either in HV or under a gas pressure. The advantage of using the gaseous environment is that positive ions, created by interactions between the incident electrons and gas molecules, neutralize the charges accumulated at the sample surface [2]. Several p-doped Si multilayers with dopant concentrations from 2·1017 atm/cm3 to 3·1019 atm/cm3, elaborated by RP-CVD, were investigated. The SIMS profile and the LVSEM image of a sample is given in Figure 1. One can clearly distinguish the difference in contrast between the five doped layers. As mentioned in the literature [1], the doping contrast is substantially affected by the surface quality (oxidizing, contamination, …) and by several experimental parameters (scan time, magnification, accelerating voltage, …). Indeed after a few minutes of investigation, we detect a decrease of the contrast. At lower magnification, the formation of a “window” can be observed in the region scanned previously (Figure 2). This phenomenon was already presented in the literature as a contamination effect [3] but it also can be due to charges accumulated at the surface of the sample. In order to have no contrast decrease and thus obtain a good reproducibility, we developed a procedure for the sample preparation and observation, which includes in situ decontamination of the sample. Using that procedure, reproducible images can be obtained and the quantification of the dopant level can be performed with a good accuracy.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 25–26, DOI: 10.1007/978-3-540-85226-1_13, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. SIMS profile and LVSEM observation on a p-doped Si multilayer structure. The SEM image was made on the Zeiss Ultra55 at 1 kV and WD = 3 mm in HV.
Figure 2. LVSEM image on the same sample at lower magnification, in HV. 1. 2. 3.
P. Kazemian, C. Rodenburg, C.J. Humphreys, Microelectronic Engineering 73-74 (2004), 948-953. A.M. Donald, Nature Materials 2 (2003), 511-516. W.H. Bruenger, H. Kleinschmidt, W. Hässler-Grohne, H. Bosse, J. Vac. Sci. Technol. B15 no. 6 (1997), 2181-2184.
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NiSi2/Si interface chemistry and epitaxial growth mode S.B. Mi1, C.L. Jia1, K. Urban1, Q.T. Zhao2, and S. Mantl2 1.Institute of Solid State Research and Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons (ER-C), Research Centre Juelich, D-52425 Juelich, Germany 2.Institute of Bio- and Nanosystems (IBN1), and Center of Nanoelectronic Systems for Information Technology (CNI), Forschungszentrum Jülich GmbH, D-52425 Jülich, Germany
[email protected] Keywords: Interface, Microstructure, Transmission Electron Microscopy
In recent years, lots of efforts have been contributed to produce epitaxial thin films of transition metal silicides on silicon substrate for fundamental studies on electrical and structural properties of interface, as well as for application as Schottky and ohmic contacts on silicon-based devices [1-3]. Since the electronic properties in metal/semiconductor junctions strongly depend on the interface structure, i.e., coordination of metal atoms, reconstruction and defects at the interface, it is of great interest to exploring the atomic structure of the interface. In this work, NiSi2 thin films are formed by annealing a Ni film deposited on a sulfur implanted-silicon (100) wafer [4]. With Surfer, which can tune the Schottky barrier height of NiSi/Si contacts [5], the interface in the NiSi2/Si system brings an additional interest of research. The atomic structure and chemistry of the NiSi2/Si interface are investigated by means of aberration-corrected transmission electron microscopy. The films grow epitaxially with an atomically sharp interface defined mainly of the (100) plane with a few {111} facets. At the (100) interface, 5-coordinated Si atoms facing 7-coordinated Ni atoms, form a 2 × 1 reconstruction structure. At the {111} facets, Ni atoms are remain to be 7-coordinated while Si atoms are 4-coordinated as in bulk. Interfacial dislocations with Burgers vectors a/4<111> and a/2<110> are observed near {111} facets with the extra half atomic planes in the Si substrate, which is the reverse of what happen between bulk NiSi2 and Si. This novel phenomenon can be understood in the light of growth modes of NiSi2 thin films with a high concentration of S at the interface. 1. 2. 3. 4. 5.
R.T. Tung，J.M. Gibson，J.M. Poate, Phys. Rev. Lett. 50 (1983) 429-432. H. von Känel, Mater. Sci. Reps 8 (1992) 193-269. C. Schwarz, H. von Känel, J. Appl. Phys. 79 (1996) 8798-8807. Q.T. Zhao, S.B. Mi, C.L. Jia, U. Breuer, S. Lenk, S. Mantl, et al., to be submitted Q.T. Zhao, U. Breuer, E. Rije, St. Lenk, S. Mant, Appl. Phys. Lett. 86 (2005) 062108.
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Figure 1 A cross-sectional low-magnification TEM image of NiSi2/Si viewed along the <011> zone axis. The (100) interface is indicated by a horizontal arrow. {111} facets are indicated by vertical arrows.
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Detailed investigation of a tunnel oxide defect in a flash memory cell using TEM-tomography U. Muehle1, M. Krause1, F. Goetze2, D. Wolf 3, and U. Gaebler4 1. Qimonda Dresden GmbH & Co. OHG, Koenigsbruecker Str. 180, D-01099 Dresden, Germany 2. Institute of Material Science, TU Dresden, D-01062 Dresden, Germany 3. Triebenberg Laboratory, ISP, TU Dresden, D-01062 Dresden, Germany 4. Infineon Dresden GmbH & Co. OHG, Koenigsbruecker Str. 180, D-01099 Dresden, Germany
[email protected] Keywords: tomography, flash memory cell, gate oxide
TEM-investigations are of increasing relevance for process monitoring and physical failure analysis in semiconductors industry [1]. However, due to decreasing dimensions of structural elements a situation can be predicted, where the details to be observed are smaller than the thickness of TEM-foil, prepared by the usual FIB-technique for target preparation [2]. Furthermore, several issues of failure analysis are limited in electrical methods for localisation, leading to an uncertainty about the optimised location of the final specimen. Therefore TEM-tomography might be able to localise and characterise the physical reason of an electrical failure by reconstruction up a 3D-model. [3] In the present case a flash memory cell showed a charge loss, probably caused by an anomaly at the active silicon surface or in the tunnel oxide. A primary investigation was performed at a TEM-foil with a thickness of more than 150 nm, including the whole width of the cell (Figure 1a). The TEM bright field image contained a suspicious detail, as shown in the close up view (Figure 1b), leading to the assumption of a so called “pit” in the surface of silicon. The comparison of its dimension to foil thickness explains the poor contrast behaviour. Further reduction of foil thickness by FIB is risky, because no information about the defect location over the foil thickness is available. For tomographic investigation a rearrangement of the sample for realisation of high tilt angles is required. In order to achieve tilt angles as high as possible, the sample was rearranged using in-situ lift-out technique [4]. The free standing sample allowed the acquisition of a tilt series from -65° till +65° using HAADF-STEM to reduce dynamic effect of the crystalline materials. The investigation was done in a Tecnai F20, equipped with a tomography samples holder type “Fischione 2020”. For data acquisition and reconstruction the FEI software-packet of Explore3D/Inspec3D was used. [5] The defect can easily be identified by slicing through the reconstructed 3D-volume. Location in any direction, lateral dimension and depth of the “pit” can be determined (Figure 2). The comparison with results of the electrical test leads to conclusions, how to avoid the this kind of defects or to guarantee a successful coverage by testing.
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Vallett, D.P.; IEEE Transact. 1 (2002); p. 117-121 http://www.future-fab.com/documents.asp?d_ID=3004 Kubel, C., Kubel, J., Kujawa, S., Luo, J.S., Lo, H.M., Russell, J.D.; 8th Int. Works. on Stress-Induced Phenomena in Metallization, AIP Conf. Proc., Vol. 817, (2006), pp. 223-228 http://www.omniprobe.com/pdf/sp1.pdf Tomography acquisition and reconstruction guide; FEI-company; Manual
Figure 1. Overview of a failing flash memory cell (a-left side) and suspicious detail with poor contrast (b-right side) in TEM-bright field.
Figure 2. Virtual slicing through the 3D-model in the plane of the “pit” in two directions.
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Overgrowth of the Mn4Si7 phase on/around the hexagonal SiC and cubic MnSi impurity phases in the Mn4Si7/Si films A. Orekhov1, T. Kamilov2 and E.I. Suvorova1,3 1. A.V. Shubnikov Institute of Crystallography of Russian Academy of Sciences, Moscow, Russia 2. Tashkent State Aeronautical Institute, Uzbekistan 3. CIME, Ecole Polytechnique Fédérale de Lausanne, Lausanne, Switzerland
[email protected] Keywords: Mn4Si7/Si films, impurity phases, TEM and X-ray EDS analysis
Structural and morphological analysis of Mn4Si7 films using the TEM/HRTEM/ Electron diffraction/STEM/EDS modes in a FEI CM300UT FEG electron microscope was carried out in a FEI CM300UT FEG electron microscope in TEM/HRTEM/ Electron diffraction/STEM/EDS mode. The images were recorded on a Gatan 797 slow scan CCD camera and processed with the Gatan Digital Micrograph 3.11.1 software, INCA (Oxford) and JEMS package [1]. The preparation of the films was performed by solid phase reaction during the evaporation of the resublimed Mn powder on the (110) Si at 950 - 1150°C and base pressure 10-4 torr. Some carbon-doped films were also grown and investigated. For optimization of growth conditions the process was carried out under equilibrium and quasi-equilibrium conditions. Obtained under quasi-equilibrium conditions the continuous Mn4Si7 film had a rough surface and a grain structure with an average size of single crystalline grains of a few microns (Figure 1ab). The carbon in the C-doped films formed a nanocrystalline SiC 100÷170 nm-layer that has divided the film into two halves (Figure 1ab). This result differs from another one previously observed for MnSiCx films [2]. The explanation can be first in different growth conditions and second that the X-ray diffraction used in [2] did not possess the capabilities for the local analysis as TEM/EDS. The Mn4Si7 islands on Si substrate were obtained under equilibrium conditions instead of continues film (Figure2a). Sizes of islands are in the range from a few nm to several tens microns. The process of reactive diffusion and formation of Mn4Si7 islands occur in the Si substrate at a depth of a few microns. Electron diffraction analysis showed that micron-size Mn4Si7 islands can contain some cubic MnSi phase inclusions. The sizes of these inclusions could reach 200-250 nm. Figure 2b shows a MnSi inclusion where [112] MnSi // [4 12 5] Mn4Si7. The electrical properties for the films were measured: resistivity and thermoelectric power were about 6⋅10-2 Ω cm and 200-250 μV/K, correspondingly 1. 2. 3.
P. Stadelmann, JEMS. http://cimewww.epfl.ch/. 2008. C. Sürgers, M. Gajdzik, G. Fischer at al. Physical Review B 68 (2003), 174423. We kindly acknowledge the help of CIME-EPFL to use the equipment and software.
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Figure 1. (a) STEM image of the continues Mn4Si7 film with the SiC layer, EDS linescan is shown for Mn, Si and C elements; (b) TEM and SAED for the single crystalline Mn4Si7 grain and nanocrystalline SiC layer.
Figure 2. (a) TEM image of Mn4Si7 micro- and nanoislands; (b) the inclusion of cubic MnSi in the Mn4Si7 microisland and the corresponding SAED pattern.
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HR-STEM EELS analysis of silicon 32 nm technology using a TITAN with a probe Cs corrector R. Pantel1, J.L Rouvière2, E. Gautier4, S. Denorme1, C. Fenouiller-Beranger3,1, F. Boeuf1, G. Bidal1 and M. Cheynet5. 1. STMicroelectronics, 850 rue Jean Monnet F-38926 Crolles, France 2. CEA DRFMC 3. CEA LETI 4. CNRS SPINTEC (2, 3, 4) address: MINATEC 17 rue des Martyrs, F-38054 Grenoble, France. 5. SIMaP (UMR 5266) CNRS-INPG-UJF, F-38402 Saint Martin-d’Hères, France
[email protected] Keywords: semiconductor silicon, EELS spectroscopy, HR-STEM, probe Cs corrector.
In this communication we present HRSTEM/EELS studies of 32 nm technology materials and interfaces using a TITAN microscope with a probe Cs corrector. Two examples are presented: silicon direct bounding and metal gates integrating HfZrO2/TiN. It is shown that EELS spectra contain rich material signatures which can help to understand transistor performance variability. The HRSTEM-EELS experiments are carried out at 300 keV using a TITAN with a CEOS probe Cs corrector installed at MINATEC. It is equipped with a Gatan TRIDIEM energy filter and two annular dark field detectors (Gatan ADF and Fischione HAADF). Figure 1a shows a HRSTEM dark field image of a direct silicon bounding (DSB) interface. The purpose of this structure is to provide two crystal orientations in order to improve the carriers mobility for both n and p MOS. The fabrication process uses Si-Si wafer bounding in presence of a thin oxide layer which is removed by high temperature annealing. The STEM Z-contrast image shows an abrupt interface and a dark area 3 nm thick in the top layer. Inside this area no oxygen is detected using EELS analysis, so this dark contrast could be due to vacancies or strain field. Figure 1 b shows a HRSTEM DF image of the bottom of a metal gate integrating Si/SiO2/HfZrO2/TiN. We have observed that on such gates the Vth and the equivalent oxide thickness (EOT) are changing continuously vs the TiN thickness (3 to 15 nm). EELS profiles presented in Figure 2 show that oxygen and nitrogen concentrations in the TiN layer depend on its thickness. In thin TiN layer, the oxygen level is high and the nitrogen low. For thick TiN layers the N/Ti ratio is about one, the oxygen been present only at the upper interface. In order to better understand the material properties low loss EELS profiles where also acquired across two stacks integrating a 12 nm and a 3 nm TiN layers. The recorded (E, x) spectrum profiles are displayed in Figure 3a and 3b respectively (intensity colour scale). These diagrams clearly visualise the different layers through their plasmon signatures. From the EELS line profile series one can extract individual spectrum inside the layers or at the interfaces. Figure 4a and 4b display the low loss spectra recorded in the middle of HfZrO2 (capped with 12 or 3nm TiN) and in the middle of both TiN layers. Significant differences in the low loss signatures are observed, especially in the TiN layers indicating nitrogen depletion in the 3 nm layer whereas a δ-TiN phase is formed in the 12 nm layer. Additional analyses will be carried out to interpret these signatures and S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 33–34, DOI: 10.1007/978-3-540-85226-1_17, © Springer-Verlag Berlin Heidelberg 2008
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extract electronics properties of the HfZrO2 dielectric layer [1-2]. These results will be presented at the conference including also near edge core loss fine structures analysis. 1. 2.
S. Pokrant, M. Cheynet, S. Jullian and R. Pantel. Ultramicroscopy 104 (2005) 233. M. Cheynet, S. Pokrant, F. Tichelaar and J.L. Rouvière, J. Appl. Phys. 101 (2007), 054101.
Figure 1: STEM Z contrast images, a) DSB silicon interface. b) Si/SiO2/HfZrO2/TiN.
Figure 2: EELS profiles across Si/SiO2/HfZrO2/TiN/Si-poly: (TiN 12, 10, 12, 5, 3 nm).
Figure 3: EELS Low Loss vs scan position across HfZrO2/TiN. a) TiN 12 nm b) 3 nm.
Figure 4: EELS spectra: a) inside HfZrO2, b) in TiN for TiN 3 nm (red) 12 nm (green).
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Tomographic analysis of a FinFET structure O. Richard, P. Van Marcke and H. Bender IMEC, Kapeldreef 75, B-3001 Leuven, Belgium
[email protected] Keywords: Electron Tomography, Fin Field Effect Transistor
In order to meet the International Technological Roadmap for Semiconductor (ITRS) goals [1] for the 32nm node and beyond new structures as the multiple-gate Silicon on insulator MOSFET (Metal Oxide Semiconductor Field Effect Transistor) with good short channel characteristics and high drive current are introduced [2]. One of the considered configurations is the FinFET (Fin Field Effect Transistor) structure consisting of single crystalline silicon fins with the gate line crossing in the orthogonal direction (Figure 1). Because the typical dimensions of these structures are smaller than the thickness of the TEM specimen, the overlap of different materials along the electron beam direction renders the interpretation of the conventional TEM images (Figure 2) difficult. Tomographic 3D reconstruction from high angular annular dark field scanning (HAADFS) TEM image series can overcome this problem, and is therefore likely the most suited technique in order to characterize such devices. TEM specimens with different geometries, i.e. cylindrical and plane parallel and with the fin parallel and perpendicular to the TEM specimens are prepared by FIB (Focus Ion Beam) milling (FEI, Strata). The acquisition of the HAADF STEM image series is performed with a FEI Tecnai (F30) FEG microscope operating at 300kV using the Explore3D software (FEI). The conventional plane parallel TEM specimens are mounted on the single tilt tomographic holder (Fischione) with the tilt axis parallel to the two perpendicular directions 3 and 1 or 3 and 2 subsequently, whereas the cylindrical specimens are fixed with the cylinder axis parallel to the tilt axis. The alignment of the HAADF STEM images and the 3D reconstruction via two schemes, weighted back projection (WBP) and simultaneous iterative reconstruction (SIRT) with 20 iterations, are performed with the Inspect3D software (FEI). Different slices from the 3D SIRT and WBP reconstructions obtained with the plane parallel sample presented in Figure 2a (tilt axis parallel with the direction 3) and with a cylindrical sample with the fin parallel to the electron beam at 0º (tilt range -70º/+70º, increment step 1º) are presented in Figure 3a-d and in Figure 3e-g, respectively. The HfO2 layer, ~ 2nm thick, exhibiting a bright contrast is detected at the top of the fin and below the polysilicon gate (arrows) and not at the fin sidewalls (Figure 3a), whereas this layer is clearly observed around the fin for the cylindrical sample (Figure 3f). The trapezoidal feature exhibiting a bright contrast on Figure 3b (arrow) comes from the high-k layer, indicating that the HfO2 layer is not properly etched in this area. Such information cannot be obtained from conventional TEM images. The bright contrast observed under the gate in the slice taken in the middle of the fin (arrow) (Figure 3c & 3e) is likely due to an artefact induced by the SIRT reconstruction scheme. This feature
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is less observed for the WBP reconstruction (Figure 3d). The nitride spacers (arrows) are clearly observed on the slice presented in Figure 3g. 1. 2. 3.
http://www.itrs.net/ H.-S.P. Wong, IBM J. Res. & Dev. 46 (2002), p. 133. The IMEC Device Implementation Project group is acknowledged for providing the sample.
Figure1. Scheme of a FinFET Figure 2. Cross-section TEM image with the fin structure with indication of the parallel (a) and perpendicular (b) to the TEM axes and plane labelling. specimen. a
b
b
c
c
d
a
e
f
f
g
g
e
Figure 3. Slices from the 3D reconstruction obtained with the fin parallel to the plane parallel specimen and with the tilt axis parallel to the direction 3: Fin in cross-section in the middle of the gate line (a). Slice parallel with the fin length at the level of the fin sidewall (b) and in the middle of the fin width with SIRT (c) and WBP (d) reconstruction scheme. Slices from the 3D reconstruction obtained with a cylindrical sample with the fin parallel to the electron beam at 0º: Slice parallel with the fin length in the middle of the fin width (e). Fin in cross-section in the middle of the gate line (f). Fin in cross-section further away from the gate (g).
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STEM EELS/EDX dopant analysis of nm-scale Si devices G. Servanton1, R. Pantel1, M. Juhel1 and F. Bertin2 1. STMicroelectronics, 850 rue Jean Monnet, F-38926 Crolles, France 2. CEA-LETI, MINATEC, 17 rue des Martyrs, F-38054 Grenoble, France
[email protected] Keywords: Silicon, Semiconductors, STEM EELS/EDX, Radiation damage, Si As/P doped
The mapping of dopant in Si devices is critical for their performance optimisation. Off-axis electron holography is a TEM based technique that is presently used to provide 2D dopant maps of semiconductor devices [1]. Due to the improvement of electron gun brightness and probe aberration correctors in TEM, the direct detection of dopant atoms is now possible. In this paper we demonstrate dopant profiling using either 200 keV low-dose EELS or 120 keV high-dose EDX. Our results are compared with SIMS. EDX and EELS experiments have been performed in STEM mode with a FEG TECNAI F20 using a EDAX detector and a Gatan Energy filter GIF 200. Samples were prepared and cleaned at low energy (5 keV) with a FEI DB400 FIB. High-energy electrons induce knock-on damage in Si [2], but below a critical energy, the momentum conservation does not allow enough energy transfer and the radiation damages become thus negligible. We evaluate in Figure 1. the critical beam energy for an intense focused probe on Si. The EDX Si signal is monitored to track the etching of the Si during electron beam irradiation. As the specimen is etched, the signal decreases and finally, a hole is observed (see inset in Figure 1.(a)). A 75-nm-thick lamella is completely etched for a (2.5 nA x 20 s) dose at 200 keV; etching also occurs at 160 keV for a dose that is 5 times higher, at 120 keV etching is negligible. This critical energy (120 keV) is lower than the quoted 200 keV [2]. Figure 1.(c) shows the etching probability and Si/Ge EDX signal measured for the same dose using SiGe reference. The 120 keV energy is then ideal for high dose EDX analysis since EDX cross-section of Si (or P) and Ge (or As) are increased. To illustrate EDX analysis versus energy, Figure 2. presents STEM EDX profiles of a P and As doped Si junction at different energy compared to SIMS profile. Each energy gives reasonable results but the dose and signal-to-noise ratio can be increased for an energy below 120 keV. Figure 3. shows the results from a 0.6% As doped Si layer analysed using low dose EELS at 200 keV and high dose EDX at 120 keV. The EELS measurements have high spatial resolution whereas the EDX profile has an excellent signal-to-noise ratio. Figure 4. shows a 120 keV EDX analysis of a BiCMOS transistor. The STEM EDX profile shows As dopant in the emitter (0.3%) and a strong segregation at the epitaxial interface (1%). This process is not optimum for transistor performances. Figure 4. shows a comparison of SIMS (acquired in 300x300 µm² box) with EELS As/Ge profiles in another BiCMOS. EELS and SIMS data are similar both quantitatively and qualitatively. In this case, the process is optimized and EELS allows studying size structure effects on epitaxial growth.
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In conclusion, the detection of dopant at concentrations 5×1019 cm-3, or below, is possible using EELS/EDX. Other applications will be presented, in particular a study of structure dimension effect during epitaxy and dopant mapping. 1. 2.
W. D. Rau, Phys. Rev. Lett. 82, 2614 - 2617 (1999) D. B. Williams, TEM, A Textbook for Materials Science, Springer (1996), p.63
Figure 1. Visualisation of Si damage using EDX signal with the effect of (a) Si specimen thickness, (b) electron beam energy, and (c) Si etching probability and Si/Ge cross-sections at different beam energies.
Figure 2. P/As junctions analysis with SIMS and EDX profiles for various energies.
Figure 3. (a) EELS (b) EDX profiles, (c) STEM image of a 0.6% As doped Si layer.
Figure 4. (a, c) BiCMOS transistors, (b) EDX, (d) SIMS and (e) EELS As/Ge profiles.
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Electron beam induced damage: An atom-by-atom investigation with TEAM0.5 C. Kisielowski1, R. Erni1, and J. Meyer2 1. National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, One Cyclotron Rd., Berkeley CA 94720, USA 2. Department of Physics, University of California at Berkeley CA 94720, USA
[email protected] Keywords: InGaN, In2O3, CdS, Ge, C, Aberration-corrected electron microscopy, single atom detection, radiation damage
A next generation electron microscopes is currently being developed by the Department of Energy as a collaborative effort to redesign the instruments around aberration corrected optics [1]. Within this project, the TEAM 0.5 prototype microscope is currently being commissioned. The instrument is equipped with a high brightness gun and a monochromator [2]. However, already in the past concerns were raised [3] and debated [4] that the high current density in field emission microscopes may alter the atomic structure of materials too fast to record undamaged images of compound semiconductors. Concerns about electron beam induced knock-on damage are indeed very relevant because the TEAM project aims at reconstructing the three dimensional structure of materials at atomic resolution, which requires maintaining structural integrity. On the other hand detailed knowledge about knock-on (and ionization) damage in such microscopes is absent and the TEAM0.5 microscope is ideally suited for such investigations since its unprecedented performance allows for the detection of single atoms of most elements of the periodic system. This contribution addresses electron beam induced damage at the single atom level in a variety of materials including InGaN alloys, In2O3, CdS, Ge and graphene. Two examples are shown in Figures 1 and 2. We find that previous investigations of beam induced atom and atom column displacements at the gold vacuum interface [5] are also observed in seminconductors and that electron beam damage scales with the strength the atoms are bound to their neighbours. Therefore, surface or edges atoms are affected most and the electron beam commonly stimulates surface diffusion. However, even if extreme illumination conditions are employed as reported in Figures 1 and 2 we find that beam induced damage is controllable at the single atom level and that even the amorphous structure of weakly bound adsorbates can be resolved in single images if 80kV of acceleration voltage is utilized. [6] 1. 2.
http://ncem.lbl.gov/TEAM-project/index.html C. Kisielowski, B. Freitag, M. Bischoff, H. van Lin, S. Lazar, G. Knippels, P. Tiemeijer, M. van der Stam, S. von Harrach, M. Stekelenburg, M. Haider, S. Uhlemann, H. Müller, P. Hartel, B. Kabius, D. Miller, I. Petrov, E. A. Olson, T. Donchev, E.A. Kenik, A. Lupini, J. Bentley, S. Pennycook, I.M. Anderson, A.M. Minor, A.K. Schmid, T. Duden, V.
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3. 4. 5. 6.
Radmilovic, Q. Ramasse, M. Watanabe, R. Erni, E.A. Stach, P. Denes, U. Dahmen, submitted (2008). M. Galtrey, R. Oliver, M. Kappers, C. Humphreys, D. Stokes, P. Clifton, and A. Cerezo, Applied Physics Letters 90 (6), 061903 (2007). C. Kisielowski and T. Bartel, Applied Physics Letters (2007). V.A. Martin, K. Ishizuka, C. Kisielowski, and L.J Allen, Phys. Rev. B 74, 172102 (2006). The TEAM project is supported by the Department of Energy, Office of Science, Basic Energy Sciences. NCEM is supported under Contract # DE-AC02-05CH11231.
Figure 1. Reconstructed phase images of In2O3 [011] from a focal series of 26 lattice images. a) Phase from images 1-13 with model inserted b) Phase from images 14-26. The crystal is oxygen terminated. The arrow marks an atom column that disappears after recording image 13. Reconstruction aperture: 0.5 Å; 300 kV; recording interval between images: 3 sec; current density: 3 104 e/Å2s; monochromator switched on.
Figure 2. a) Lattice image of graphene with monolayer adsorbates. Carbon atoms are imaged bright. b) The lattice fringes of the monolayer graphene sheet are removed from the image. The amorphous structure of the adsorbate is resolved in single images. 80 kV; current density: 1.4 105 e/Å2s; monochromator switched on.
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The atomic structure of GaN-based quantum wells and interfaces C.J. Humphreys1, M.J. Galtrey1, R.A. Oliver1, M.J. Kappers1, D. Zhu1, C. McAleese1, N.K. van der Laak1, D.M. Graham2, P. Dawson2, A Cerezo3 and P.H. Clifton4. 1. Department of Materials Science and Metallurgy, University of Cambridge, Pembroke Street, Cambridge CB2 3QZ, UK. 2. School of Physics and Astronomy, Alan Turing Building, University of Manchester, Manchester M13 9PL, UK 3. Department of Materials, University of Oxford, Parks Road, Oxford OX1 3PH, UK 4. Imago Scientific Instruments, 5500 Nobel Drive, Madison, Wisconsin 53711, USA
[email protected] Keywords: gallium nitride, quantum wells, interfaces
We have used a combination of high resolution electron microscopy (HREM), three dimensional atom probe (3DAP) microscopy and atomic force microscopy (AFM) to reveal the atomic structure of InGaN quantum wells (QWs) and InGaN interfaces. Such quantum wells and interfaces are of considerable scientific and technological importance because they form the basis of GaN-based LEDs and lasers. A long-standing puzzle has been why GaN-based LEDs emit brilliant light when the dislocation density is so high. Originally it was believed that this was because the excitons (bound electron-hole pairs which recombine, emitting light) were localized in indium-rich clusters in the InGaN quantum wells and so could not diffuse to dislocations, which would quench the light emission. However, careful HREM work showed this was not the case [1], and this was confirmed by 3DAP images that showed that InGaN was a random alloy, up to at least 25% In content [2]. If there are no In-rich clusters in the InGaN QWs, what mechanism is responsible for localizing the excitons and preventing them from diffusing to dislocations? HREM images suggest that the lower GaN/InGaN QW interface is flat and the upper InGaN/GaN QW interface has monolayer steps several nm across [3]. 3DAP images strikingly confirm this, see Figure 1 [4]. Because of the strain in the QW and the strong piezoelectric effect in InGaN, a monolayer interfacial step is sufficient to localize the carriers at room temperature and prevent them from diffusing to dislocations [3]. If an InGaN QW is grown at the same temperature as the GaN barriers, a uniform QW results (with interface steps on its upper surface). However, two-temperature growth is commonly used, in which the GaN barriers are grown at a higher temperature than the InGaN QW. In this case HREM, AFM and 3DAP reveal that the quantum well can break-up into a network structure, on a 50-100nm lateral length scale. The threading dislocations mostly pass through the gaps in the QWs, and this provides an additional localization mechanism, on a longer length scale, see Figure 2 [5,6]. 3DAP reveals that a low level of In is present in the GaN barriers although the In flux is switched off during growth of the barriers, and this occurs for both blue and
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green emitting QW structures. The source of this In appears to be the InGaN QW layers, with In segregating, not diffusing, into the barriers from the QWs [4]. We conclude that HREM, 3DAP and AFM form a powerful combination of microscopy techniques for studying the atomic structure of quantum wells and interfaces in semiconductors. 1. 2. 3. 4. 5. 6.
T. M. Smeeton, M. J. Kappers, J. S. Barnard, M. E. Vickers and C. J. Humphreys, Appl. Phys. Lett. 83 (2003), 5419. M. J. Galtrey, R. A. Oliver, M. J. Kappers, C. J. Humphreys, D. J. Stokes, P. H. Clifton and A. Cerezo, Appl. Phys. Lett. 90 (2007), 061903. D. M. Graham, A. Soltani-Vala, P. Dawson, M. J. Godfrey, T. M. Smeeton, J. S. Barnard, M. J. Kappers, C. J. Humphreys and E. J. Thrush, J. Appl. Phys. 97 (2005), 103508. M. J. Galtrey, R. A. Oliver, M. J. Kappers, C. J. Humphreys, P. H. Clifton, D. Larson, D. W. Saxey and A. Cerezo, J. Appl. Phys. (2008), submitted. N. K. van der Laak, R. A. Oliver, M. J. Kappers and C. J. Humphreys, Appl. Phys. Lett. 90 (2007), 121911. M. J. Galtrey, R. A. Oliver, M. J. Kappers, C. McAleese, D. Zhu, C. J. Humphreys, P. H. Clifton, D. Larson and A. Cerezo, Appl. Phys. Lett. 92 (2008), 041904.
Figure 1. 3D atom probe image of the (a) upper and (b) lower interfaces of an InGaN/GaN quantum well emitting in the green (25% In content in the well), showing interface steps on the upper interface.
50 nm
Figure 2. Cross-sectional STEM-HAADF image of a commercial bright green LED showing gross thickness variations (arrowed) in all four InGaN quantum wells.
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Using TEM to investigate antiphase disorder in GaP films grown on Si(001) T.B. Adams1, I. Nemeth1, G. Lukin1, B. Kunert2, W. Stolz1, K. Volz1 1. Materials Science Centre, Philipps-University Marburg, D-35032, Germany 2. NAsP III/V GmbH, D-35041, Germany
[email protected] Keywords: dilute nitride; defects; MOVPE
GaP-based dilute nitride materials are of great technological interest since Ga(NAsP) exhibits a direct band-gap and can be grown pseudomorphically-strained on Si substrates. Multi-quantum-well laser devices grown on GaP substrates have already been demonstrated with efficient and stable performance[1], and the growth of such devices on nominally exact Si(001) would represent a tremendous technological leap forward, allowing the incorporation of optoelectronic devices into silicon chips by conventional CMOS fabrication. To date, however, completely defect-free dilute nitride films grown directly on Si(001) have not been achieved. Since Ga(NAsP) can be grown defect-free on GaP, and GaP can in principle be grown on Si(001), one possible solution is to use GaP as a buffer layer between the Si substrate and Ga(NAsP) layer(s). Growth of defect-free GaP on Si presents its own challenges: conditions must be such that planar defects such as twins and stacking faults are eliminated. Anti-phase domain boundaries (APB) represent a special case of planar defect, where the polarity of the crystal is inverted on either side of the domain. Polarity inversion can be regarded as a 90o rotation of the crystal about the axis perpendicular to the plane of the substrate/film interface. APB’s originate from inhomogenous nucleation of the 1st monolayer of GaP, from mono-steps on the Si surface, or from a thermodynamic need to distribute interfacial charge at the GaP/Si interface. The density of APB nucleation sites on exact-oriented Si substrates can be minimised by appropriate substrate preparation and the remaining APB’s can, under certain growth conditions, ‘kink’ so that they connect and self-annihilate before reaching the GaP surface. APB’s can therefore be tolerated in the GaP buffer layer, provided they have self-annihilated before reaching the buffer layer surface. The TEM is an indispensable tool for characterising APB morphology and for providing critical information to allow optimisation of growth conditions. APB’s can be imaged directly using HR-TEM and dark field imaging (see Figure 1, for example), while the polarity can be determined locally using convergent beam electron diffraction. The application of these techniques in the growth of GaP on Si(001) by Metal Organic Vapour Phase Epitaxy will be discussed and defect structure will be correlated to growth conditions. 1.
B. Kunert, A. Klehr, S. Rheinhard, K. Volz, W. Stolz, Electronics Letters 42 (2006)
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Figure 1. Dark field image pair showing morphology of antiphase domains in a GaP buffer layer grown in Si(001).
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TEM characterization of InAs/GaAs quantum dots capped by a GaSb/GaAs layer A.M. Beltrán1, T. Ben1, A.M. Sánchez1, D.L. Sales1, M.F. Chisholm2, M. Varela2, S.J. Pennycook2, P.L. Galindo3, J.M. Ripalda4 and S.I. Molina1 1. Dep. Ciencia de los Materiales e Ing. Met. y Qca. Inorg., Facultad de Ciencias, Universidad de Cádiz. Campus Río San Pedro, Puerto Real 11510 Cádiz, Spain 2. Materials Science and Technology Division, Oak Ridge National Laboratory, 37831 Oak Ridge, Tennessee, USA 3. Dep. Lenguaje y Sistemas informáticos, CASEM, Universidad de Cádiz, Campus Río San Pedro, Puerto Real 11510 Cádiz, Spain 4. Instituto de Microelectrónica de Madrid (CNM-CSIC) Isaac Newton 8, 28760 Tres Cantos, Madrid, Spain
[email protected] Keywords: InAs QDs, TEM, GaSb
It is well known that there is intense interest in expanding the usable wavelength for electronic devices. Research efforts are dedicated to develop GaAs technology in order to achieve emission at 1.3 and 1.55 µm, so GaAs could be used for telecommunication applications [1]. Ga(As)Sb on InAs/GaAs quantum dots (QDs) is a promising nanostructure to be used in telecommunications. The introduction of antimony to these nanostructures is an effective solution to obtain redshift emission, even at room temperature [2]. Several authors have reported a miscibility gap for GaSb and InAs in absence of strain [3]. Our preliminary studies by Conventional Transmission Electron Microscopy (CTEM) of nanostructures on GaAs (001) substrates show the formation of a quaternary alloy InxGa1-xAsySb1-y when a GaSb layer is grown over the QDs. In order to achieve a proper formation of the system InAs QDs/GaSb barrier layer, the growth of an intermediate GaAs layer between the InAs QDs and the GaSb cap layer resulted as a promising growth solution. In this work we have studied the effect of the introduction of intermediate layers with diverse thickness comparing with InAs QDs directly capped. Studied samples have been grown by molecular beam epitaxy on GaAs (001) substrate. InAs was deposited, followed by a GaSb or GaAs/GaSb layers. Finally, the nanostructure was capped by GaAs [4]. In this work we have analysed sample A (without GaAs intermediate layer), sample B1, B2 and B3 with 3, 6 and 12 MLs of GaAs intermediate layer respectively. The structural characterisation of these samples was carried out by TEM in a JEOL 1200EX and 2011EX at 120 and 200 kV respectively. During this study dark field (DF) g002 cross-section (XTEM) images were taken from samples A, B1 and B2 (see Figure 1a for sample A). These images show in general dots with a brighter core than the GaAs substrate which are formed on and surrounded by GaAs. In these images the GaSb layer can not be clearly distinguished. In accordance with g002DF intensities dynamically simulated by EMS software (not shown in this work), GaSb layers have brighter
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contrast than GaAs and InAs layers for sample thickness lower than 50 nm (case conscientious chosen for this study). In view of these results a GaSb layer mixed with GaAs, and maybe with InAs (from QDs and intermediate layers), seems to have been formed. However, on sample B3 (see Figure 1b) a layer with brighter contrast than GaAs substrate can be undoubtedly observed on the intermediate layer, which can be related to a GaSb layer. It could be said that depositing 12 MLs of GaAs avoids the intermixing of GaSb with GaAs intermediate layer and, therefore, with InAs. In order to clarify the type of InGaAsSb alloy inside dots, compositional studies initially on sample A were carried out by high resolution scanning transmission electron microscopy (STEM) using a high-angle annular dark field (HAADF) detector and spatially resolved electron energy loss spectroscopy (EELS) in two dedicated scanning transmission microscopes. Taking GaAs as reference for Ga and As, In and Sb maps were calculated as 100-CGa and 100-CAs, respectively (CGa and CAs represent Ga and As concentration) from low loss EEL spectra. These studies have demonstrated the presence of the four elements inside the QDs in an estimated concentration of 40% (~20% at the wetting layer). These results have been confirmed by the analysis of HAADF images. Since in this technique higher intensity is associated with higher atomic numbers (Z), the position of the maximum of an intensity profile would correspond with the area of the highest concentration of In/Sb. The maximum is located in the central region of the QD which is associated with the quaternary alloy InxGa1xAsySb1-y [5]. In conclusion, it has been shown that the presence of a GaAs intermediate layer between the InAs QDs and the GaSb layer improves structural quality and minimise the intermixing of the binary compound semiconductor constituting the nanostructure, as the GaAs thickness increases over 6 MLs. Compositional analyses have demonstrated the formation of a quaternary alloy inside the QDs when GaSb is grown over InAs QDs. 1. V.M. Ustinov and A.E. Zhukov, Semicond. Sci. Technol. 15 (2000) p. R41. 2. J. M. Ripalda, D. Granados, Y. González, A.M. Sánchez, S.I. Molina and J.M. García, Appl. 3. 4. 5. 6.
Phys. Lett. 87 (2005) 202108. G.B. Stringfellow, J. Cryst. Growth 58 (1982) p. 194. J. M. Ripalda, D. Alonso-Álvarez, B. Alén, A. G. Taboada, J. M. García, Y. González and L. González, Appl. Phys. Lett. 91 (2007) 012111. S.I. Molina, A.M Sánchez, A.M Beltrán, D.L Sales, T. Ben, M.F. Chisholm, M. Varela, S.J Pennycook , P.L. Galindo, A.J. Papworth, P.J. Goodhew and J.M Ripalda, Appl. Phys. Lett. 91 (2007) 263105. This work has been supported by MEC (TEC2005-05781-C03-01 and 02/MIC), SANDiE European network (Contract NMP4-CT-2004-500101), Junta de Andalucía (PAI research group TEP-120) and U.S. Department of Energy, Division of Materials Science and Engineering (MFC, MV and SJP). We acknowledge Prof. P. J. Goodhew and Dr. A. J. Papworth for the use of the STEM facilities in the University of Liverpool. a)
QDs
GaAs substrate
GaAs cap layer
25 nm
b) GaAs substrate
GaAs cap layer QD
GaSb
Figure 1. XTEM DF g002 images of samples a) A and b) B3.
25 nm
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The microstructure of (0001)GaN films grown by molecular beam epitaxy from a nanocolumn precursor layer D. Cherns1, L. Meshi1, I. Griffiths1, S. Khongphetsak1, S.V. Novikov2, N. Farley2, R.P. Campion2, and C.T. Foxon2 1. Department of Physics, University of Bristol, Tyndall Avenue, Bristol, UK 2. Department of Physics and Astronomy, University of Nottingham, Nottingham, UK
[email protected] Keywords: GaN, polarity, crystal growth, dislocation reduction
A major challenge in the production of GaN devices is to reduce the densities of threading dislocations, typically up to 109 – 1011 cm-2, generated by the use of highly mismatched substrates such as (0001)sapphire. One method is to use epitaxial lateral overgrowth (ELO) where a mask is used to restrict the growth of the GaN to “seed” columns. Lateral growth leads to “wings” where dislocation densities may be several orders of magnitude less than in the seed. In this paper we consider a maskless ELO approach in which the GaN is grown on (0001)sapphire. Firstly, GaN is grown by molecular beam epitaxy under strongly N-rich conditions, leading to pronounced nanocolumn growth [1]. The growth is then continued under Ga-rich conditions, leading to lateral growth and coalescence of the nanocolumns. Eventually, continuous GaN layers are formed with threading dislocation densities in the range 108 – 109 cm-2. Growth was carried out using plasma-assisted MBE (PA-MBE) in a Varian ModGen II MBE system. A 5nm AlN buffer layer was grown at a temperature of 600°C prior to the GaN growth at 700-800°C. Following up to 6 hours of growth in a high overpressure of active nitrogen provided by an HD25 RF activated plasma source, the N flow was reduced and growth continued for 1-6 hours under Ga-rich conditions. TEM samples were prepared in cross-sectional orientation, either by mechanical polishing followed by Ar-ion thinning in a Gatan PIPS, or by sectioning in an FEI 201 focused ion beam thinner (FIB), and examined at 200kV in a JEOL 2010 TEM. SEM studies were also carried out on as-grown samples using a JEOL JSM7400F FEGSEM. Figure 1 shows a sample grown only under N-rich conditions. This contains a rough, but mostly continuous, intermediate GaN layer extending up to 400 nm from the substrate. In addition, there are discrete GaN nanocolumns which protrude through this layer up to around 2 μm from the substrate. Close inspection indicates that the nanocolumns contain relatively few defects, with contrast mainly consisting of bend contours. In contrast, the intermediate layer is highly defective, containing both threading dislocations and high densities of basal plane stacking faults, which are near edge-on and mostly out of contrast in this figure. The polarity of the deposit was examined by taking convergent beam electron diffraction (CBED) patterns under symmetrical (0002) systematic row conditions, i.e. with ±0002 reflections at the same deviation parameter. Under these conditions, the intensities in ±0002 reflections differ owing to double
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diffraction contributions [2]. The nanocolumns were found to be Ga-polar, and the intermediate layer N-polar. Figure 3 illustrates a layer grown firstly under N-rich conditions and then under Garich conditions. The nanocolumns have grown laterally near the column tops, leading to a continuous overlayer. Coalescence of the nanocolumns gives appreciable areas which are free of threading defects, but with occasional threading defects describing the coalescence of nanocolumns which are locally misoriented. In contrast, the N-polar underlayer has undergone negligible further growth under Ga-rich conditions. Figure 3 demonstrates that nanocolumn growth provides a successful method of reducing threading dislocation densities in GaN films. In the paper, the mechanisms of nanocolumn growth and of dislocation reduction will be illustrated and explained. 1. 2. 3.
E. Calleja et al , Phys. Stat. Sol. b 244, 2816-2837 (2007) D. Cherns et al, Phil. Mag.A77 273-286 (1998) This work was supported under an EPSRC grant EP/D080762/1
Figure 1. SEM (left) and TEM (right) showing Ga-polar nanocolumns protruding from an intermediate layer of N-polar GaN, grown for 6 hrs by MBE under N-rich conditions.
Figure 2. GaN layer grown for 5 hrs under N-rich conditions and then 5 hrs under Garich conditions showing a continuous Ga-polar overlayer (growth surface: bottom right).
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Compositional and Morphological Variation in GaN/AlN/AlGaN Heterostructures P.D. Cherns, C. McAleese, M.J. Kappers and C.J. Humphreys Department of Materials Science and Metallurgy, University of Cambridge, Pembroke St, Cambridge, CB2 3QZ, UK
[email protected] Keywords: AlGaN, STEM, EFTEM
AlGaN/GaN heterostructures are an important materials system for the development of optoelectronic devices emitting in the ultraviolet region of the electromagnetic spectrum. Due to a difference in lattice parameter between GaN (a=0.31893 nm) and AlN (a=0.31130 nm), AlGaN grows under tensile strain on GaN. This leads to cracking above a critical thickness, which decreases with increasing Al content. It has been shown that this problem can be overcome by the use of low [1] or high [2] temperature-grown AlN interlayers. The function of the interlayer is to relax sufficiently without macroscopic cracking, in order to exert compressive strain on the subsequently grown AlGaN layers. In this paper, we present a study of AlN interlayers by high-angle annular dark field (HAADF) imaging and energy-filtered transmission electron microscopy (EFTEM). The work focuses on the morphological and compositional variations that result from changes in growth temperature and growth time. The samples are grown on c-plane sapphire by metal organic vapour phase epitaxy (MOVPE). Interlayer growth temperatures were varied between 700°C and 1020°C, and thicknesses between 1 nm and 50 nm. It is shown that the use of AlN interlayers thicker than 5 nm is effective in preventing cracking in AlGaN layers, which is independent of growth temperature. However, it has been seen that changes in growth temperature do have a significant impact on the composition and morphology of the interlayers. High resolution HAADF images of 5 nm and 10 nm interlayers grown at 1020°C are presented in Figure 1. The images show examples of frequently observed trench-like features in the interlayers, as well as an indication of changes in interlayer composition. To study composition changes at high and low growth temperature, EFTEM filtered series were acquired around the Ga L2,3 edge at 1115 eV. A 20 eV energy selecting slit was used and 18 images were acquired at 20 eV spacing. The resultant Ga maps, as presented in Figure 2, shows that at higher growth temperatures the Al content is inhomogeneous in the interlayers. Ga is present in the interlayer in contrast to the pure AlN that is expected. There is vertical segregation of material, with Al-rich material near the top of the layer. The trenches are observed more frequently as the growth temperature is increased and experimental evidence indicates that Al is favourably incorporated around the edges of the features. Possible mechanisms for the formation of these trenches are discussed.
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This study has important implications for the understanding of the growth of AlGaN/GaN heterostructures, particularly in relation to the incorporation of Al under tensile strain. 1. 2.
H. Amano, M. Iwaya, N. Hayashi, T. Kashima, S. Nitta, C. Wetzel and I. Akasaki, Physica Status Solidi (b) 216 (1999), p683. I.H. Lee, T.G. Kim and Y.Park, Journal of Crystal Growth 234 (2002), p305.
Figure 1. HR-HAADF images from samples with (a) 5 nm and (b) 10 nm AlN interlayers, grown at 1020°C. Images are acquired with sample oriented along a <1120> zone axis.
Figure 2. (a) EFTEM Ga maps, presented both raw and corrected for thickness variation and elastic scattering, taken from samples with 30 nm AlN interlayers grown at 700°C and 1020°C. (b) Intensity profiles taken from the two corrected Ga maps, normalised to 100% Ga content in the GaN template.
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TEM/STEM/EFTEM imaging and Valence Electrons Spectroscopy analysis of Ultra Low-K dielectrics M. Cheynet1, S. Pokrant2, M. Aimadeddine3, V. Arnal3, F. Volpi1 1. Laboratoire SIMaP 1130 rue de la Piscine BP75 Saint Martin d’Hères 38042 France 2 Carl Zeiss NTS GmbH, 73447 Oberkochen, Deutschland. 3. ST Microelectronics Rue Jean Monnet, 38920 Crolles, France
[email protected] Keywords: Ultra-Low-k, electronic properties, TEM-STEM imaging, EFTEM, VEELS.
The ITRS requires the integration of dielectric materials with effective dielectric constant (k) lower than 2.8 [1]. This is achieved using porous organic silicate glass in full-sheet deposition. Unfortunately, during the integration processes several steps like lines etch or Cu Chemical Mechanical Polishing (CMP) introduce damages [1]. One issue, to avoid this damage, is the deposition of a SiO2 based capping layer before the CMP. In this work, we report on the results deduced from imaging techniques performed in the environment of TEM (i.e. BF/DF TEM and STEM imaging, EFTEM and Z-contrast imaging), and on the results deduced from Valence Electron Energy Loss Spectroscopy (VEELS) analysis, both performed to quantify the capping layer effect on the density, the chemistry and the dielectric properties of the low-k layer. 1 – Microstructure and density of “capped” and direct CMP low-k layers: STEM bright/dark field and EFTEM images are recorded from 100 nm thick fibbed TEM foils of “capped” and “direct” CMP low-k samples. For “direct” CMP samples, severe damages are revealed within the Low-K layer. These damages are observed at depth of several tens of nanometers from the polishing surface. They mainly correspond to a decrease of the Low-K material density as indicated by the zero and plasmon loss profiles recorded across the Low-K layers. On the contrary, for “capped” samples, homogeneous contrast and profile are observed in the whole Low-K layer thickness, indicating that, neither significant physical damages nor chemical changes occur. No severe chemical damage is observed, except a small carbon depleted layer wrapping the Cu interconnects lines as shown in the EFTEM images displayed in Figure 2. 2 – Band gap mapping across the “capped” and “direct” CMP ULK layers. SiOCH based low-k samples are transparent to electron for a thickness higher than 200 nm. This allows using thicker TEM foils than usually for low-loss analysis and excludes surface excitation modes to be detected in the low-loss spectra. Moreover, since the refractive index of conventional Low-K materials are reported to be lower than 1.4, Cerenkov losses can be also excluded even working at 200 kV acceleration voltage. The signatures observed in the single scattering low-loss spectra (SSD) can be thus all attributed to electronic transitions after the removal of the zero-loss peak and of the plural scattering from the experimental spectrum. Similar SSD are obtained independently of ZL peak deconvolution procedure applied (Pearson VII function, Richardson Lucy deconvolution [2]). Figure 3 illustrates the band gap determination for
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the CMP capped sample. Figure 4 illustrated the result of the band gap and defects energy mapping across the capped and direct CMP Low-K layers. The nanometer scale low-loss maps show that the band gap profile within the capped ULK layer is flat, while it is deeply disturbed in the “direct CMP” layer. The CMP process generates defects close to the conduction band, which is detrimental for the electrical behaviour of the integrated ULK layers. The protection role of the SiO2 based capping layer is clearly demonstrated. 1. K. Maex, M. R. Baklanov, D. Shamiryan, F. Lacopi, S.H. Brongersma and Z. S. Yanovitskaya, J. Appl. Phys. 93, 8793 (2003).
2.
Acknowledgement to Jaysen Nelayah from the LPS d’Orsay in France for the treatment of the low-loss spectra using a Richardson-Lucy deconvolution
Figure 1: TEM cross-section image of the 5th ML of the studied stack, for the capped (left) and the direct CMP (right) samples. Zero-loss and low-loss profiles recorded in the Low-K layer showing strong density variations in the direct CMP sample. Figure 2: C-K edge, O-K edge, C-K/O-K edges EFTEM images showing a carbon depleted layer wrapping the Cu lines
Figure 3. Band gap and defects energy transition determination for “capped” low-k layer.
Figure 4. Band gap and defect energy signature profiles along the ULK layer of the “direct CMP” sample (left) and the “capped” sample (right) showing the protection role of the SiO2 based capping layer
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Epitaxial Orientations of GaN Grown on R-plane Sapphire J. Smalc-Koziorowska1,3, G.P. Dimitrakopulos1, Ph. Komninou1, Th. Kehagias1, S.-L. Sahonta1, G. Tsiakatouras2 and A. Georgakilas2 1. Department of Physics, Aristotle University of Thessaloniki, GR 54124 Thessaloniki, Greece 2. Microelectronics Research Group, Department of Physics, University of Crete, P.O. Box 2208, 71003 Heraklion-Crete, Greece, and IESL, FORTH, P.O. Box 1527, GR 71110 Heraklion-Crete, Greece 3. Department of Materials Science, Warsaw University of Technology, 02-507 Warsaw, Poland
[email protected] Keywords: nonpolar Gallium Nitride, Transmission Electron Microscopy, epitaxy
Nitride-based heterostructure nanostructures grown along non-polar plane orientations are recently investigated for optoelectronic device applications. Such orientations are promising in order to eliminate the high interface charge densities and spatial separation of carrier wave functions in quantum well structures, which are due to the built-in polarization along the [0001] direction of the wurtzite structure. One possible substrate for growth of a-plane ( 1 1 20 ) III-nitrides is r-plane ( 1 1 02 ) sapphire, and growth of both GaN and AlN epilayers on such substrates by various epitaxial techniques has already been reported. In this work the studied gallium nitride films were grown by nitrogen radio-frequency plasma source molecular beam epitaxy (RFMBE) on nitridated r-plane sapphire substrates. Transmission electron microscope (TEM) investigation has shown the presence of misoriented gallium nitride islands at the epilayer/substrate interface, with average diameter ~10 nm “Figure 1”. The misoriented crystallites were overgrown by the nonpolar a-plane GaN, and they were promoted by the presence of interfacial steps on the sapphire surface, introduced by the nitridation. The matrix aplane GaN exhibited the [ 1 100 ]GaN//[ 1 1 20 ]sapphire [0001]GaN//[ 1 101 ]sapphire epitaxial relationship, whereas the orientation relationship of the misoriented islands was [ 01 1 0 ]GaN//[ 1 101 ]sapphire , (0001)GaN~//( 2 1 1 3 )sapphire “Figure 2(a)”, i.e. the islands were semipolar. It was determined that the islands nucleated directly on the sapphire surface and in particular on { 2 1 1 3 } step risers. Furthermore, pockets of sphalerite structure were introduced at low growth temperature as a result of the superposition of basal-plane stacking faults along the [0001] growth direction inside these grains “Figure 2(b)”. Crystallographically equivalent variants of the secondary orientation led to the introduction of twinning “Figure 2(a)”. These phenomena were usually confined close to the epilayer/substrate interface, although some larger misoriented islands reaching the epilayer surface were also observed. [1] 1.
Work supported under the European Community's Marie Curie Research Training Network PARSEM (MRTN-CT-2004-005583)
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Figure 1. Cross-section xTEM bright-field image taken under two-beam conditions off the [ 1 100 ]GaN zone axis with gGaN = 0002. The sapphire substrate is out of contrast. Areas of moiré fringes indicate the presence of misoriented islands.
(a)
(b)
Figure 2. (a) Selected area electron diffraction pattern including the sapphire substrate (black solid line, [ 1 101 ]sapph.za), nonpolar a-plane GaN (white solid line, [0001] za) and two variants of semipolar GaN (black and white dashed lines, < 01 1 0 >GaN za). (b) Cross-section HRTEM image of the interfacial area between GaN epilayer and sapphire substrate viewed along the [ 1 011 ]sapphire direction inclined at ~77o relative to the substrate normal.
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Microstructure and growth model of MBE-grown InAlN thin films S.-L. Sahonta1, A. Adikimenakis2, G.P. Dimitrakopulos 1, Ph. Komninou1, H. Kirmse3, E. Pavlidou1, E. Iliopoulos2, A. Georgakilas2, W. Neumann3 and Th. Karakostas1 1. Solid State Physics Section, Department of Physics, Aristotle University of Thessaloniki, 54124 Thessaloniki, Greece 2. Microelectronics Research Group, IESL, FORTH, PO Box 1527, 71110 Heraklion, Crete, Greece and Department of Physics, University of Crete, PO Box 2208, 71003 Heraklion, Crete, Greece 3. Institut für Physik, Humboldt-Universität zu Berlin, AG Kristallographie, Newtonstrasse 15, D-12489 Berlin, Germany
[email protected] Keywords: indium, nitrides, GaN, MBE, HEMT
III-nitride materials are used in a variety of optoelectronic devices due to their high carrier mobility and high current densities. Recently optoelectronic waveguides and field effect transistor junctions are being fabricated with a new type of heterostructure based on In0.18Al0.82N/GaN, as this provides a large band gap difference at the heterojunction coupled with lattice-matching to GaN, inhibiting the formation of straininduced defects known to reduce carrier recombination and mobility [1]. Despite its attractive properties, InAlN growth is challenging owing to the very different optimum growth temperatures for InN and AlN films, requiring InAlN to be grown at a temperature low enough to prevent dissociation of InN, yet high enough to allow sufficient aluminium adatom surface diffusion to deposit high crystalline quality film. There are to date very few studies on the microstructure and morphology of InAlN films, yet such observations are vital in order to fabricate InAlN-based devices with the greatest performance, and to understand the basic structural properties of InAlN. In this work transmission electron microscopy (TEM) is used to characterise a variety of high quality InAlN films grown by molecular beam epitaxy (MBE) on AlN/sapphire (0001) and GaN/sapphire (0001) substrates. Indium alloy contents are varied to allow different conditions of epitaxial lattice mismatch. For all mismatch conditions studied it is found that the InAlN films are deposited as hexagonal columnar grains with distinct {11-20}- and {1-100}-faceted grain boundaries, often lined with threading dislocations. A corresponding transition from 2D to 3D growth mode is observed by reflection high energy electron diffraction (RHEED) during the deposition of the first few nanometres of film, indicating island coalescence. Crystallite grain size is on average 15 nm in diameter regardless of InAlN interfacial lattice mismatch, with slight lateral misalignment of adjacent crystals, resulting in a rough stepped film surface, confirmed by atomic force microscopy (AFM). TEM mass contrast and highangle annular dark field (HAADF) studies (Figure 1) show indium accumulation at the grain boundaries in between the columns. Energy-dispersive X-ray spectroscopy
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(EDXS) shows an In content of 51 ± 2 atomic % at the grain boundaries. Eventual film planarisation and compositional homogeneity does not occur after prolonged growth, with the In-rich grain boundaries continuing to the film surface even in thick films (Figure 2). A new growth model is proposed in which the observed non-uniform indium distribution is attributed to film deposition kinetics. 1.
J. Xie, X. Ni, M. Wu, J.H. Leach, Ü. Özgür and H. Morkoç, Applied Physics Letters 91 (2007), p. 132116.
2.
Support by the European Community's Marie Curie Research Training Network PARSEM (MRTN-CT-2004-005583) is gratefully acknowledged.
(a)
(c)
(b)
10 nm
10 nm
10 nm
Figure 1 (a). [0001] TEM image of In0.19Al0.81N film on GaN. Grains of around 15 nm diameter have facets aligned with {1-100} and {11-20}. Pronounced dark mass contrast at grain boundaries indicates accumulation of heavy In atoms; bright mass contrast InAlN shows high Al content in grain centres. (b) [0001] TEM image of In0.12Al0.88N film on AlN shows the same structure as (a). (c) HAADF image of the same film as (b), heavy In atoms appear with bright AlN contrast, confirming segregation of In to the grain edges. Figure 2. Cross sectional bright-field TEM image of a thick In0.12Al0.88N film on GaN, taken with g = 0002, showing periodic vertical contrast due to threading dislocations and scattering from vertical grain boundaries.
100 nm
2
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Silicon carbide modulated structures as a result of the introduction of 8H bands in a 4H matrix N. Frangis1, M. Marinova1, I. Tsiaoussis1, E.K. Polychroniadis1, T. Robert2, S. Juillaguet2, J. Camassel2 1. Solid State Physics Section, Department of Physics, Aristotle University of Thessaloniki, GR-54124 Thessaloniki, Greece 2. Groupe d'Etude des Semiconducteurs, CNRS and Université Montpellier 2, cc 074-GES, 34095 Montpellier cedex 5, France
[email protected] Keywords: silicon carbide, high-resolution electron microscopy, modulated structures
Due to its excellent physical properties, silicon carbide (SiC) is so far considered as a promising wide band gap semiconducting material for high temperature, high frequency and high power electronic devices. In the present work we report on the structural characterisation of a 4H-SiC epitaxial layer grown by Chemical Vapour Deposition (CVD) in a horizontal, resistively heated, hot wall CVD reactor [1]. Stacking faults (SFs) are very common defects in all SiC polytypes, because of their small formation energy. Conventional transmission electron microscopy found in some areas of the studied material a very high density of defects, parallel to the basal hexagonal planes. High-resolution transmission electron microscopy (HRTEM) revealed that no isolated SFs appear. Instead the observed defects are three-dimensional ones, having the form of bands with a stacking sequence different than that of the 4H structure (…ABCBABCB…). These bands can be considered as the result of repeated SFs and (or) twin planes, within a distance of a few layers. Two types of bands were found very often to occur (“Figure 1a”): 1) An eight-bilayers (Si and C) band having the structure of a single unit cell of the 8H polytype, which has a stacking sequence: ABCABACB. 2) A fourteen-bilayers band. It contains three sub-bands, with four, six and four bilayers, respectively. Each sub-band has the cubic stacking sequence and is separated from the others by twin planes. In the electron diffraction patterns, taken from the highly defected regions of the specimens, satellite spots appear between the 4H reflections. Two sequences of spots were observed (“Figure 2”), denoting that at least two incommensurate modulated structures are locally created. HRTEM images reveal that the origin of these modulations is the quasi-periodic introduction of the defect bands. In the micrograph of “Figure 1b” a series of 8H single unit cells are seen, separated by a number (not always an integer one) of 4H unit cells. 1. 2.
T. Chassagne, A. Leycuras, C. Balloud, P. Arcade, H. Peyre and S. Juillaguet, Mat. Science Forum 457-460 (2004), p. 273. We kindly acknowledge that this work was performed in the frame of the MANSIC project (MRTN-CT-2006-035735), financed by EU.
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Figure 1. a) A HRTEM micrograph showing the presence of 8H and 14TC (14 layers twinned cubic) bands, having a different stacking sequence than the 4H matrix. b) A HRTEM micrograph, where the introduction of the 8H single unit cell bands occurs in a quasi-periodic way, giving rise to the appearance of satellite spots in the electron diffraction patterns.
Figure 2. Central [000l]* rows from electron diffraction patterns, denoting the local formation of SiC incommensurate modulated structures. The 4H reflections are indicated.
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HRTEM study of AlN/3C-SiC heterointerfaces grown on Si(001) and Si(211) substrates T. Isshiki1, K. Nishio1, Y. Abe2, J. Komiyama2, S. Suzuki2, H. Nakanishi2 1. Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606-8585, Japan 2. Covalent Materials Corporation, 30, Soya, Hadano, Kanagawa 257-8566, Japan
[email protected] Keywords: AlN, SiC buffer layer, Interface structure, HRTEM
Epitaxial growth of nitride semiconductors on Si wafer is an important technique to develop electro-optical devices. However, a large lattice mismatch between Si and nitrides disturbs direct growth of defect-free nitride layer on Si substrate. Cubic SiC (3C-SiC), which can be formed directly on Si substrate [1], is attractive as a buffer layer to relieve the lattice mismatch. Nanoscopic information at interface between nitride and 3C-SiC is required to improve quality of growing nitride layer. In this work, interface structure between AlN and 3C-SiC buffer layer grown on Si(001) and Si(211) substrates is studied by high-resolution transmission electron microscopy (HRTEM). Specimen for cross-sectional HRTEM observation was prepared as following procedure. 3C-SiC buffer layer was grown on Si(001) and (211) substrates in a low pressure VPE reactor. After carbonization of Si substrates with C3H8 at 1200 °C, 3CSiC layer was grown with SiH4 and C3H8 at 1250 °C and 6.7 kPa in H2 atmosphere. AlN layer about 50-100 nm in thickness was deposited on the 3C-SiC at 1200 °C and 3.3 kPa in a low pressure MOVPE reactor using (CH3)3Al and NH3 with H2 carrier gas. On Si(001) substrate, 3C-SiC buffer layer grew with identical crystal orientation to the substrate with SiC(001) surface. On this surface, hexagonal AlN (h-AlN) didn’t grow directly, but cubic AlN(c-AlN) grew with pyramidal shape at the AlN/SiC interface within 10 nm in thickness, as shown in Figure 1. Since lattice mismatch between SiC(001) and c-AlN (001) is as small as 1.3%, c-AlN grew at first on the SiC(001) surface. However, due to phase instability of c-AlN, h-AlN grew on the four pyramidal c-AlN{111} planes immediately with h-AlN(0001) parallel to c-AlN{111}. As a result, growth layer was composed of h-AlN domains being in twin relation. A Si(211) substrate is restructured spontaneously to undulant structure of Si(111) and Si(200) facets during carbonized process [2]. Therefore, nano-facet structure of flat SiC(111) and unclear SiC(200) also appeared at surface of growing 3C-SiC buffer layer, as shown in Figure 2. On the flat SiC(111) facet, h-AlN grew coincidently with (0001) plane parallel to the facet. On the other hand, small voids and (11, 01) twin crystal were observed near the SiC(200) facet. Stacking faults propagating from the SiC layer were absorbed by the voids and didn’t propagate in h-AlN layer. Generation of new defect at the interface is hardly observed except twin boundary mentioned above. 1.
S. Nishino, H. Suhara, H. Ono and H. Matsunami, J. Appl. Phys. 61 (1987) 4889
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T. Nishiguchi, Y. Mukai, M. Nakamura, K. Nishio, T. Isshiki, S. Ohshima and S. Nishino, Materials Science Forum 457-460 (2004) 285.
Figure 1. HRTEM image at interface of AlN grown on SiC(001) / Si(001).
Figure 2. HRTEM image at interface of AlN grown on SiC(111) and SiC(200) / Si(211).
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Electron microscopy of GaAs/AlGaAs quantum cascade laser A. Łaszcz, J. Ratajczak, A. Czerwinski, K. Kosiel, J. Kubacka-Traczyk, J. Muszalski, M. Bugajski and J. Kątcki Institute of Electron Technology, Al. Lotników 32/46, 02-668 Warsaw, Poland
[email protected] Keywords: transmission electron microscopy, quantum cascade laser, GaAs, AlGaAs
The quantum cascade laser (QCL) is based on a different principle comparing to bipolar semiconductor lasers, because it uses only one type of the charge carriers – electrons and therefore is called an unipolar laser. The essence of the QCL is an operation like an electronic waterfall. Electrons cascade down a series of identical energy steps built into the material during crystal growth, emitting a photon at each step as opposed to diode lasers which emits only one photon over the similar cycle. The other aspect of the QCL is the possibility of emission over a wide range of wavelengths using the same combination of materials in the active region [1]. The use of GaAs/AlGaAs heterostructures for preparation of unipolar laser has been demonstrated by Sirtori et al. in 1998 [2] and it was not restricted to one material system. We report electron microscopy study of GaAs/AlGaAs QCL structure fabricated by the Molecular Beam Epitaxy (MBE) technique. The QCL was designed according to the scheme described in Table I. The studies focused on MBE technology verification and precision of realization of layer’s thicknesses in the active region of the QCL. Transmission electron microscopy (TEM) observation showed a repeatability of modules in the active region (30x16 GaAs/AlGaAs layers, Figure 1a). However, an increase of each module thickness with reference to the initial assumption (Table I) has been observed. Intended thickness of one module was 45.3 nm. The one determined by TEM measurements equals 48.9 nm (Figure 1b). The ratio of measured to desired thickness of layers is 1.079. TEM results were confronted with X-ray diffraction (XRD) analysis of layer thickness in the active region. XRD results showed that the ratio of measured to desired thickness of layers is 1.075 [3]. The difference between thickness determined by both methods is only 0.4%. 1. 2. 3. 4.
J. Faist, F. Capasso, D.L. Sivco, C. Sirtori, A.L. Hutchinson and A.Y. Cho, Science 264 (1994), p. 553. C. Sirtori, P. Kruck, S. Barbieri, P. Collot, J. Nagle, M. Beck, J. Faist and U. Oesterle, Applied Physic Letters 73, No. 24 (1998), p. 3486. K. Kosiel, M. Bugajski, J. Muszalski, J. Kubacka-Traczyk, J. Ratajczak, A. Łaszcz, P. Romanowski, R. Mogilinski, J. Gaca and M. Wójcik, Design and MBE Growth, Optimization of Mid-Infrared Quantum Cascade Laser Structures - Int. Conf. on Semiconductor Materials and Optics (SMMO 2007).
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Table I. Scheme of GaAs/AlGaAs quantum cascade laser.
Figure 1. TEM images of: (a) active region of quantum cascade laser (30x16 GaAs/AlGaAs layers) and (b) high magnification of the first module of the active region with 16 GaAs/AlGaAs layers.
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Solving the crystal structure of highly disordered Sn3P4 by HRTEM O.I. Lebedev, A.V. Olenev, G. Van Tendeloo EMAT, University of Antwerp, Groenenborgerlaan 171, B2020 Antwerpen, Belgium
[email protected] Keywords: crystal structure, disordering, HRTEM
Disordered narrow-gap semiconductors have recently been investigated in the context of a search for new thermoelectric materials. Sn3P4 is a narrow-gap semiconductor exhibiting a relatively high thermal conductivity of about 8 W m–1 K–1. Upon cooling to 150 K it undergoes a remarkable transition from n-type conduction to p-type conduction. The crystal structure of Sn3P4 long remained unsolved since its discovery in 1920. In order to check possible superstructure formation Sn3P4 has been investigated by TEM. ED patterns for Sn3P4 along the main zones are shown in Fig.1. Remarkably, only the brighter spots in the [0001] ED pattern (Fig.1a) can be indexed using the X-ray single crystal data. On the [–1100] and [01–10] ED patterns (Fig.1b and c, respectively) streaks along the c* axis are clearly present. Along the [0001] zone these streaks intersect the plane and produce sharp spots (Fig. 1a, d) (marked by squares and triangles). To understand the nature of the streaks HRTEM was performed along [01-10] (Fig.2). The streaks are also clearly observed on the FT, indicating that the origin of the streaks is captured in the image. The image itself has a number of peculiarities: the rows of dots along the a axis are arranged in such a way that a bright dot is followed by a dim dot. The X-ray data unambiguously showed that while the sites of the tin atoms are fully occupied, the phosphorus atoms, joined into dumbbells, are disordered in such a way that only one out of four dumbbells can be present in each block of the structure. Consequently, the contrast modulation, observed in Fig.2a is most probably associated with a short range ordering of P atoms along the a axis. Good agreement between the experimental and the calculated image may be achieved by introducing a partially ordered model (shown in Fig.2c), in which disordered P24– dumbbells alternate along the a direction with partially ordered fragments. In these fragments, only two out of four possible dumbbells remain intact, the third is broken leaving a separate phosphorus atom, and the fourth is missing. One can see (Fig.2a marked by white frames) that modulated rows of dots in some cases are shifted along the c direction. This shift is random, which means that disordered and partially ordered phosphorus fragments irregularly alternate along the c axis as shown in Fig.2c. This irregularity leads to the formation of streaks in the [–1100] and [01–10] ED patterns. An image of the basic Sn3P4 structure can be obtained by applying a Bragg-mask filter for only the basic spots (correspond to the brighter reflections on the ED pattern). Inverse FT results in the image shown in Fig. 2b. In this case, a uniform intensity distribution in the rows of bright dots is achieved. An image calculated on the base of the X-ray structure S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 63–64, DOI: 10.1007/978-3-540-85226-1_32, © Springer-Verlag Berlin Heidelberg 2008
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refinement (inset Fig. 2b) perfectly fits and shows that the rows of bright dots correspond to the positions of Sn(1) in the average crystal structure.
Figure 1. ED patterns along the major zones of Sn3P4
Figure 2. (a) [01-10] HRTEM image of Sn3P4 and its corresponding FT pattern. The SnP block stacking along the c axis is marked with white quadrangles. The modulation along the Sn layer is marked by white arrows; (b) filtered HRTEM image of (a). Only the basic spots were included in the mask; (c) proposed structure model of the disordered structure. Simulated images based on a disordered model (c) and the ordered model (X-ray data) are given as inset in (a) and (b).
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Determination of In-distribution in InGaAs quantum dots D. Litvinov1, H. Blank1, R. Schneider1, D. Gerthsen1 and M. Hetterich2 1. Laboratorium für Elektronenmikroskopie, Universität Karlsruhe, D-76128 Karlsruhe 2. Institut für Angewandte Physik, Universität Karlsruhe, D-76128 Karlsruhe
[email protected] Keywords: quantum dot, InGaAs, TEM
We investigated InAs/GaAs quantum dot (QD) layers by quantitative transmission electron microscopy (TEM) and photoluminescence (PL) spectroscopy. To achieve the highest possible In-concentration in the QDs and maximum QD sizes, InAs QD layers with a nominal thickness of 2.6 ML were grown by molecular-beam epitaxy on GaAs (001) with a low deposition rate of 0.006 monolayers (ML)/sec at 500 °C. To reduce the strong In-segregation, which is observed in InAs/GaAs heterostructures during the growth of the GaAs cap layer, the InAs QD layers were capped by InGaAs with a nominal thickness of 6 nm and finally by a GaAs cap layer with 30 nm thickness. The influence of different In-concentrations x in the InxGa1-xAs cap layer (x = 0, 0.05, 0.1, 0.15, 0.2 and 0.25) on the structural and optical properties was studied. The Indistribution in the QD layers was determined with the composition evaluation by lattice fringe analysis technique (CELFA) [1] which is based on the chemical sensitivity of the (002) reflection for the sphalerite structure. The measured In-concentration in the QDs is artificially lowered due to the averaging effect of a TEM sample with finite thickness and the three-dimensional nature of the QDs embedded in a matrix with lower Inconcentration. A post-processing procedure was developed and applied to the raw CELFA data to determine the real In-concentration in the QDs. We combine conventional, quantitative high-resolution and high-angle annular darkfield scanning transmission electron microscopy of plan-view and cross-sectional samples to determine the shape and size of the quantum dots. The shape of QDs, which is a truncated pyramid with square bases along the <100> directions and with {101} facets, is essentially independent of the cap-layer composition. To determine the TEM specimen thickness in the QD region, a tilt series of cross-section images was taken with the chemically sensitive (200) reflection. For the CELFA evaluation, only QDs were chosen which are fully embedded in the Ga(In)As cap layer, i.e. without any sectioning due to TEM specimen preparation. For the post-processing we calculate a function which - applied to the CELFA data multiplies values of the local In-concentration by the quotient of the local thickness of the sample and the local thickness of the QD. Figures. 1a,c are color-coded maps of the In-concentration obtained with CELFA for the samples capped by GaAs and In0.2Ga0.8As. For comparison, Figures. 1b,d display the corresponding maps after the post-processing with a other color scale. The thicknesses of the TEM specimens according to the tilt series are 33 nm for the GaAs-capped QD and 38 nm for the In0.2Ga0.8As-capped QD. Since the correction was only applied to the QD region, the wetting layer is not visible in the corrected maps. The imaging was carried out with the
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(200) lattice fringes (perpendicular to the substrate/layer interface), which are less affected by lattice plane bending in contrast to the (002) lattice planes parallel to the growth direction. The white triangles indicate the approximate QD shape and size. The lateral size of the QDs increases with the In-concentration in the cap layer from about 20 nm for the sample with x = 0 to 24 nm for the island with x = 0.2, while the QD height increases only slightly from 6.8 to 7.3 nm. The white arrows mark the maximum of the In-concentration within the QDs which shifts upward after the correction. The determination of the composition of the layers after post-processing shows that the Inconcentration in the QDs increases in growth direction and reaches values up to 90 %. Redistribution of indium during the InGaAs cap layer growth leads to a decrease of the In-concentration in the cap layer with respect to the nominal In-concentration. The observed red shift of the PL peak with increasing In-concentration in the cap layer is attributed to the enlargement of island size and the change of the strain in the QD layers caused by the misfit between InAs and the cap layer. 1. 2.
A. Rosenauer, Transmission Electron Microscopy of Semiconductor Nanostructures - An Analysis of Composition and Strain (Heidelberg, Springer) Springer Tracts in Modern Physics (2003), 182 p. This work has been performed within the project A.2 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. Colour-coded In-concentration maps obtained with CELFA using the (200) reflection in samples capped by GaAs (a,b) and In0.2Ga0.8As (c,d) before (left-hand side) and after post-processing of the CELFA data (right-hand side). The white triangles delineate the QD shape.
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Nanoanalytical investigation of the dielectric gate stack for the realisation of III-V MOSFET devices. P. Longo1, A.J. Craven1, M.C. Holland2 and I.G. Thayne2 1. Department of Physics and Astronomy, University of Glasgow, G12 8QQ, UK 2. Department of Electrical & Electronics Engineering, University of Glasgow, G12 8LT, UK
[email protected] Keywords: TEM, EELS, STEM, compositional analysis, MOSFET, gallium arsenide
Planar Si MOSFET technology using Si(ON) is rapidly approaching its theoretical limit and the search for new material is essential. n-type GaAs has a mobility 5 times higher than Si [1]. This makes it a suitable candidate for MOSFETs devices. So far, GaAs has not been used for practical MOSFETs because of the difficulties of making a good dielectric oxide layer in terms of leakage current and unpinned Fermi Level. Using processes pioneered by Passlack et al. [2], dielectric gate stacks consisting of a template layer of amorphous Ga2O3 followed by amorphous GdGaO have been grown on GaAs substrates. Careful deposition of Ga2O3 can leave the Fermi Level unpinned. The introduction of Gd is important in order to decrease the leakage of current. Electron energy-loss spectroscopy (EELS), carried out in a nano-analytical electron microscope, has been shown to be an excellent method to characterise Ga2O3/GdGaO dielectric stacks [3]. Spectrum imaging (SI) uses computer control to position the electron beam and to record one or more spectra at each point, building up an Ndimensional data set. It is also possible to quantify the amount of each element present, with extremely high spatial resolution. Such approach has been applied to the characterization of a number of GaAs/Ga2O3/GdxGa0.4-xO0.6 systems where x is the Gd atomic fraction. The thickness of the GdGaO layer varies from 9nm up to 60nm. In general, the Gd concentration decreases with the distance from the substrate. However we have observed an increase in one case where the peak Gd concentration was 36%. For Gd target concentration of 20%, the films are normally amorphous. As the Gd concentration increases, crystallisation can start to occur at >25%. However we have also observed the start of crystallisation at lower Gd concentration under specific growth conditions. “Figure 1” shows a high-resolution TEM bright field image of the GaAs/Ga2O3/GdGaO dielectric gate stack where the highest Gd content detected is 36%. The high Gd concentration causes the partial crystallisation of the GdGaO layer as shown in Figure 1. The Fourier Transform inset in “Figure 1” also shows the preferred orientation of the lattice structure in the GdGaO layer. XRD analysis has confirmed the presence of crystalline Gd2O3 in the GdGaO layer. The avoiding crystallisation is an important factor during the growth of the dielectric stack for MOSFET devices. The optimum conditions in terms of device properties are when the GGO layer is completely amorphous. “Figure 2” shows the elemental composition across the GdGaO layer. The
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thickness profile ranges from 0.22 down to 0.1t/λ. The Gd concentration slightly increases on moving away from the Ga2O3 layer and reaches the highest value in the area at 22nm. On the hand the O concentration appears to be slightly below 60%. However in the region in the top surface of the GdGaO layer the O concentration increases significantly. A sharp drop of t/λ is seen in the region where the O concentration increases. Thus the increase in the O concentration is likely to be due to the presence of surface layers with modified composition whose effects become stronger as the sample gets very thin. 1. Wang YC, Hong M., Kuo JM., Kwo J., Mannaerts JP, Chen YK and Cho Y., IEEE, Electron Device Letters 20, 457, 1999 2. Passlack M, Yu Z, Droopad R, Bowers B, Overgaard C, Abrokwah J, and Kummel AC, J Vacuum Science & Technology B17, 1, 49-52, 1999 3. Longo P, Craven AJ, Scott J, Holland MC and Thayne IG proceedings of MSM XV, 2007 4. Authors would like to acknowledge EPSRC support under grant EP/F002610 and Mr B. Miller for TEM specimen preparation.
GdGaO
Ga2O 3 GaAs Figure 1. High-resolution TEM bright field image of the GaAs/Ga2O 3/GdGaO dielectric gate stack. The GdGaO layer is partly crystalline with preferential orientation, as shown in the inset Fourier Transform. 0.3 0.25
50
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40 30
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20
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Gd% O% t/lambda
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Figure 2. Elemental concentrations across the GdGaO layer. The 0 on the x-axis represents the main boundary between the Ga2O 3 template layer and the GdGaO. The thickness profile is also shown.
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Phase mapping of uncapped InN quantum dots J.G. Lozano1, M. Herrera2, R. García1, N.D. Browning2,3, S. Ruffenach4, O. Briot4 and D. González1 1.Departamento de Ciencia de los Materiales e Ingeniería Metalúrgica y Química Inorgánica, Universidad de Cádiz, Apdo 40, E-11510 Puerto Real (Cádiz), Spain 2. Department of Chemical Engineering and Materials Science, University of California-Davis, One Shields Ave, Davis, CA 95616, USA 3. Chemistry, Materials Science and Life Sciences Directorate, Lawrence Livermore National Laboratory, 7000 East Avenue, Livermore, CA 94550, USA 4. Groupe d’Etudes des Semiconducteurs, UMR 5650 CNRS, Place Eugène Bataillon, Université Montpellier II, 34095 Montpellier, France
[email protected] Keywords: Indium Nitride, Indium Oxide, HRTEM, Phase mapping
InN has in recent times attracted much attention due to its outstanding electronic properties and recently set low band gap of 0.7 eV [1]. However, it has been the nitride that has generated higher controversies about its fundamental parameters, due to the difficulties associated to its growth and to the application of standard measurement techniques developed for other III-V materials. Many of these variations in the measured values are ascribed to oxygen contamination [2], either during the growth process, either as a simple natural aging process. The aim of this work is to obtain maps of the location of the different phases inside the InN nanostructures, by using high resolution transmission electron microscopy (HRTEM) combined with the geometric phase (GP) algorithm [3]. The samples were grown by metalorganic vapour phase epitaxy [4] onto sapphire substrates, using a thick (1 μm) GaN buffer layer grown at a temperature of 1000º C; on top of which InN quantum dots were grown at a temperature of 550º C. The experimental analysis were carried out in a JEOL 2010FEG and JEOL 2500 transmission electron microscopes operating at 200 kV. Figure 1(a) shows a cross sectional HRTEM of an InN QD, and Figure 2(b) its associated Fast Fourier Transform (FFT). In the last, several phases can be distinguised: wurtzite InN oriented along the <11-20> zone axis, two bcc-In2O3 crystals oriented along the <110> pole and rotated 180º around the growth direction one with respect to each other and a InN zincblende phase, whose main spots share position with the In2O3 phases. To reconstruct the crystal position we used the GP algorithm, and the results are shown in Figure 2. It can be seen that the hexagonal InN phase has disappeared in the areas near the surface shrinking to the centre of the QD. Additionally, the two sets of inclined fringes correspond cubic phases (zinc-blende InN and bcc-In2O3). Therefore, in the final configuration cubic phases form a surrounding layer that envelops a core of InN-w. Although In2O3 constitutes the mainly component of the cubic regions, the presence of InN-zb is also reported
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1. 2. 3. 4.
A.G. Bhuiyan, A. Hashimoto, A. Yamamoto, J. Appl. Phys. 94, 2779 (2003) M. Yoshimoto, H. Yamamoto, W. Huang, H. Harima, J. Saraie, A. Chayahara and Y. Horino Appl. Phys. Lett., 83, 3480 (2003) O. Briot. B. Maleyre, S. Ruffenach, Appl. Phys. Lett. 83, 14 (2004) M.J. Hÿtch, E. Snoeck and R. Kilaas, Ultramicroscopy 74 (1998) 131
Figure 1.(a) X-HRTEM micrograph of an uncapped InN QD taken along the <11-20> pole. In the inset, FFT of this micrograph (b) Diffraction patterns simulated over the FFT image where yellow triangles correspond to InN-w, white and pink squares correspond to both InN-zb grains and red and green circles simulate the In2O3 grains.
Figure 2. Numerical moiré fringes mappingof the uncapped QD taking as reference the spot g11-2 spot of the In2O3 crystals. The radius of the Bragg mask was chosen to include the {222} spots of the FFT images (see blue circle in Fig.1b)
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STEM investigations of (In,Ga)N/GaN quantum wells P. Manolaki1, I. Häusler1, H. Kirmse1, W. Neumann1, P. Vennéguès2, P. De Mierry2 and P. Demolon2 1. Humboldt-Universität zu Berlin, Institut für Physik, AG Kristallographie, Newtonstrasse 15, D-12489 Berlin, Germany. 2. Centre de Recherche sur l'Hétéro-Epitaxie et ses Applications, rue Bernard Grégory, 06560 Valbonne, France
[email protected] Keywords: Quantum wells, (In,Ga)N, STEM
Nanostructures of III-N semiconductors have a great commercial potential due to their optoelectronic properties. The structural and chemical characteristics of such structures determine the performance of the light-emitting devices. Transmission electron microscopy can be used to determine these characteristics at a nanometer scale. Nanostructures were generated by metal organic chemical vapour deposition on sapphire (0001). First a GaN buffer layer doped with Si was grown on the substrate. In a second step 5 (In,Ga)N/GaN multilayers of a nominal In content of 8.5 at. % were grown. The nominal thickness of the (In,Ga)N layers was 2.6 nm separated by a 22 nm thick GaN barrier layer. The structure was capped by a 200 nm thick GaN layer doped with Mg. For the p-contact a double layer consisting of 5 nm Ni and 5 nm Au was deposited. The (In,Ga)N quantum wells (QWs) were grown at a temperature of 755 °C, whereas the GaN barriers were grown at 975 °C. The sample was prepared for TEM by mechanical preparation comprising formatting, mechanical polishing, dimpling, and finally ion milling. Structural and chemical analysis was performed on a TEM/STEM JEOL JEM2200FS operating at 200 kV. The instrument is equipped with a high-angle annular dark-field (HAADF) detector and an energy dispersive X-ray (EDX) spectrometer. The minimum probe size in STEM mode is 0.14 nm. In the composition-sensitive STEM HAADF image of the Figure 1 the 5 QWs as well as the Ni/Au contact layer appear brighter than the substrate, the buffer layers and the cap layer, because they contain elements of a higher atomic number. In addition, Figure 1 shows a number of screw and edge dislocations originating from the buffer. Figure 2, which is a magnification of the marked area of Figure 1, evidences that the individual (In,Ga)N layers exhibit inhomogeneous intensity as well as varied thickness. Figure 3 is a STEM image of another area of the sample, where just below the contact layer trapezoidal pits are visible. However, they are not initiated at threading dislocations. The results of the analysis of the chemical composition of the (In,Ga)N/GaN QWs and of the barriers by EDXS show an In content of 4 to 8 at% which is below the nominal value of 8.5 at% (see Figure 4). The results of the EDX line scan along the first QW are given in Figure 5, verifying a compositional inhomogeneity. The In composition for the QWs can be also evaluated by using strain analysis on high-resolution TEM images of the (In,Ga)N QWs as described in [1]. S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 71–72, DOI: 10.1007/978-3-540-85226-1_36, © Springer-Verlag Berlin Heidelberg 2008
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1. 2.
P. Ruterana, S. Kret and M. A. Poisson, Mater. Sci. Eng. B 93, (2002), p. 185. This work was financially supported by the Marie Curie Research Training Network PARSEM (contract No.: MRTN-CT-2004-005583).
Figure 1. STEM HAADF overview of the QW, of the contact layer and of dislocations, respectively
Figure 2. Magnification of the framed area of Figure 1
Figure 4. EDX spectrum of the line scanning across the layers (see green arrow on Figure 2)
Figure 3. STEM overview of the QWs and the contact layer of another area of the sample
Figure 5. EDX spectrum of the line scanning along the first QW (see blue arrow on Figure 2)
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Defects in m-plane GaN layers grown on (100) γ-LiAlO2 A. Mogilatenko1, W. Neumann1, T. Wernicke2, E. Richter2, M. Weyers2, B. Velickov3 and R. Uecker3 1. Institut für Physik, Humboldt Universität zu Berlin, Newtonstr. 15, 12489 Berlin, Germany 2. Ferdinand-Braun-Institut für Höchstfrequenztechnik, Gustav-Kirchhoff-Str. 4, 12489 Berlin, Germany 3. Institut für Kristallzüchtung, Max-Born-Str. 2, 12489 Berlin, Germany
[email protected] Keywords: GaN, thin films, crystal defects
In the last years (100) γ-LiAlO2 was used as an alternative substrate for GaN epitaxy [1, 2]. In contrast to the commonly used commercial substrates (100)LiAlO2 allows a nearly lattice-matched growth of both polar c-plane and non-polar m-plane oriented GaN layers. The m-plane oriented GaN layers show absence of internal polarisation fields along the growth direction [1-1.0]GaN which is advantageous for fabrication of optoelectronic devices with high performance [2]. Furthermore, LiAlO2 decomposition at elevated substrate temperatures [3] leads to spontaneous separation of the substrate from thick GaN layers during post-growth cooling down [1, 4]. This allows the preparation of freestanding GaN layers, which can be used as substrates for subsequent GaN homoepitaxy. In the present study we applied transmission electron microscopy (TEM) to investigate the structure of m-plane GaN layers grown on LiAlO2 by metal organic vapour phase epitaxy (MOVPE) process. Plan-view diffraction contrast analysis shows the presence of different types of defects. The m-plane GaN layers contain a large number of basal plane stacking faults (BSF) which are visible in the bright-field image obtained with a (10-10)GaN reflection (Figure 1a). The nature of the stacking faults was determined in cross-section (Figure 2a). According to high resolution TEM analysis they are intrinsic stacking faults of I1 type (Figure 2b). This defect corresponds to the following stacking sequence: …ABABABCBCBC…, where each capital letter represents a layer consisting of a Ga-N atom pair. This stacking fault is characterized by a displacement vector R = {1/3[1-1.0] + ½[00.1]}. Additionally, the layers contain a complicated net of boundaries which are visible in the bright-field image obtained with a (0002)GaN reflection (Figure 1b). According to the thickness fringes in Figure 1b it seems that this net of boundaries consists of planar defects located on a number of GaN planes, which are inclined to the growth direction. These defects are characterised by a displacement vector R arranged parallel to [00.1]GaN. Cross-sectional analysis proved the presence of grain boundaries at {10-10}, {1-102} as well as {1-103}GaN planes. The inclined boundaries act as additional nucleation sites for the BSF formation (Figure 2c).
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1. 2. 3. 4.
H.P. Maruska, D.W. Hill, M.C. Chou, J.J. Gallagher, B.H. Chai, Opto-electronics Rev. 11 (2003), p. 7 P. Waltereit, O. Brandt, A. Trampert, H.T. Grahn, J. Menniger, M. Ramsteiner, M. Reiche, K.H. Ploog, Nature 406 (2000), p. 865 A. Mogilatenko, W. Neumann, E. Richter, M. Weyers, B. Velickov, R. Uecker, J. Appl. Phys. 102 (2007), p. 023519 E. Richter, Ch. Hennig, U. Zeimer, M. Weyers, G. Tränkle, P. Reiche, S. Ganschow, R. Uecker, K. Peters, Phys. Stat. Sol. (c) 3 (2006), p. 1439
Figure 1. Plan-view bright-field images of m-plane GaN layers deposited by MOVPE. The images were obtained under two-beam conditions with (a) (10-10)GaN reflection and (b) (0002)GaN reflection. Basal plane stacking faults and a net of grain boundaries are visible in (a) and (b), correspondingly.
Figure 2. (a) Cross-sectional bright-field image of m-plane GaN layer shows the presence of BSF and defects which are inclined to the [1-1.0]GaN growth direction. (b) High-resolution TEM image of a BSF proves the layer sequence:…ABABCBC… (c) A zigzag shaped grain boundary in m-plane GaN acts as a nucleation site for the BSF formation.
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Improvements on InP epilayers by the use of monoatomic hydrogen during epitaxial growth and successive annealing F.M. Morales1, A. Aouni2, R. García1, P.A. Postigo3, C.G. Fonstad4, S.I. Molina1 1. Departamento de Ciencia de los Materiales e Ingeniería Metalúrgica y Química Inorgánica, Universidad de Cádiz, 11510 Puerto Real, Cádiz, Spain 2. Faculté de’Sciences et Techniques, Université Abdelmalek Essadi, BP 416 Tanger, Morocco 3. Instituto de Microelectrónica de Madrid, Centro Nacional de Microelectrónica, CSIC, Isaac Newton 8, PTM Tres Cantos, 28760, Madrid, Spain 4. Department of Electrical Engineering and Computer Science, and Microsystems Technology Laboratory, MIT, Cambridge, Massachusetts 02139, USA
[email protected] Keywords: InP on GaAs, MBE, TEM
InP has demonstrated advantages for optoelectronics and wireless applications but there is still a lack of products which offer big InP surfaces for its microelectronics integration. The heteroepitaxy of InP films on cheap and large GaAs substrates would be an appropriate solution for this necessity since bulk InP wafers are available only in small sizes and at high prices [1]. In this way, the challenge for InP growers focuses on finding the best alternatives for avoiding the negative effects produced by the lattice mismatches and the differences in the thermal expansion coefficients between active InP films and substrate materials. Beyond the use of buffer intermediate layers between InP and GaAs, the technique of direct growth of InP on GaAs using monoatomic hydrogen has recently demonstrated excellent results. Compared to homoepitaxial InP layers grown under the same conditions, H-assisted heteroepitaxial InP showed enhanced optical properties. Hydrogen helped to a raise on the critical thickness of 40 times larger than that expected for such a high misfit on InP/GaAs [2]. The characterization of high-quality InP epilayers fabricated by low temperature (200-480ºC) solid-source molecular beam epitaxy on (001) GaAs is reported in this work. The optimized two-step production process is firstly assisted by monoatomic hydrogen in the earliest stages of the molecular beam epitaxy growth. After growth, some samples have been rapid-thermal annealed (RTA). Selected area electron diffraction (SAED) studies confirm that the InP layers are always well heteroepitaxially placed as indicated by perfect alignments between GaAs and InP diffraction lattice spots. It is concluded by the analyses of the diffraction patterns and the study of interfaces in highresolution transmission electron (HRTEM) micrographs that the zinc-blende InP is almost strain relaxed. These images allowed observing misfit dislocations located just at the interface and formed after the critical thickness was reached during the growth. Moreover, two other main kind of structural defects were observed from diffractioncontrast transmission electron micrographs (DC-TEM), both typical of cubic crystals: stacking faults (SF) and tangled threading dislocations (TDs). These imperfections are
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mostly confined to a region of less than 500 nm above the InP/GaAs interface (due to interfacial strain) and their densities decrease monotonically as they approach the surface of thee ~2 μm thick InP films (auto-annihilation and decreased stress). The TDs appear irregularly spread across the f and are revealed with g vectors parallel to <220> or <400> directions on cross-sections. However, planar defects did not appear in the samples with higher H doses and the interface region looks cleaner of imperfections. As a rule, it is found a diminution on the density of TDs on the layers when either H or RTA is used (Fig. 1). For improved samples (H+RTA), the TD density was measured to be on the range of 108 cm-2, which is a good result compared to state of the art InP epilayers [3]. However, comparable quantities of such line defects have demonstrated to improve electrical and optical properties by acting as a gettering network located beneath the device zone on InP/GaAs based high-electron mobility transistors [4]. 1 2 3 4
R. Szweda, III-Vs Review, 13, 55 (2000) P. A. Postigo, F. Suarez, A. Sanz-Hervás, J. Sangrador and C. G. Fonstad, J. Appl. Phys. 103, 013508 (2008). K.-F. Yarn, C.-I Liao and C.-L. Lin, Crystallographic Reviews, 12, 47–80 (2006). Y. Liu and H. Wang J. Appl. Phys. 100, 034505(2006)
Figure 1 HRTEM micrographs of the InP layer showing a SF (a) and a structure clean of defects (b). TDs at the surface of the optimum sample (i. e. higher content of H during growth and subsequent RTA) (c). SAED pattern showing both GaAs and InP aligned related spots without contrasts associated to defects (d). CDTEM micrographs (g=002) of the InP layers for conditions: No H and No RTA (e); High H and No RTA (f); Medium H + RTA (g) and; High H + RTA (h).
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Study of microstructure and strain relaxation on thin InXGa1-xN epilayers with medium and high In contents F.M. Morales1, J.G. Lozano1, R. García1, V. Lebedev2,3, S. Hauguth-Frank2, V. Cimalla2,3, O. Ambacher3, D. González1 1. Departamento de Ciencia de los Materiales e Ingeniería Metalúrgica y Química Inorgánica, Facultad de Ciencias, Universidad de Cádiz, 11510 Puerto Real Cádiz, Spain 2. Institute of Micro- and Nanotechnologies, Technical University Ilmenau, 98693 Ilmenau, Germany 3. Fraunhofer Institute for Applied Solid State Physics, Tullastrasse 72, D-79108 Freiburg, Germany
[email protected] Keywords: InGaN, High resolution TEM, Strain
GaN and InGaN with a high content of Ga became the most important semiconductors for light emission and detection in the visible and the near UV spectral regions. On the other hand, the fabrication of InGaN epilayers with medium and high concentration of In is still not well developed. However, the control of the growth of alloys within this In-rich range would give rise to expand the optical activity of InGaN till the near IR. Recently, the compositional dependencies of photoconductivity and electron transport properties for metal-semiconductor-metal photodetectors based on Inrich InGaN ultra-thin films were studied [1]. There, the electron density profiles and low-field mobilities for different compositions of InGaN were calculated. It was demonstrated that in contrast to bulk InN exhibiting a dominating surface electron accumulation, the free electrons in ultra-thin InxGa1-xN/GaN (0.5<x<1) heterostructures tend to accumulate mostly at the buried InGaN/GaN interface. In this work we report on the fabrication and studies of high structural quality wurtzite InGaN nanolayers grown on top of Al2O3/AlN/GaN epitaxial substrates within this unexplored range of compositions of medium and high In concentrations (0.4<x<0.8). The epilayers are single-crystalline and did not show appreciable phase segregation or compositional modulation. Electron and X-ray diffraction were used for an accurate estimation of the actual a- and c-lattice parameters. The thin InGaN layers present a tetragonal distortion (i. e. strained by biaxial compressive stress) which decrease monotonically as the In content is higher for the same film thickness. This effect was also confirmed by the observation of high-resolution transmission electron micrographs. Therefore, the accurate composition of InGaN alloys should be assessed not only by diffraction methods, but by chemical analysis to serve as an independent reference, since for biaxial compressed InGaN, Vegard’s law interpolations lead to systematic overestimations of In compositions. For this reason, the technique of energy dispersive X-ray spectroscopy in scanning transmission electron microscopy mode was carried out for the best measurement of In/Ga compositions, in the same regions, where
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selected area electron diffraction patterns were registered. Once correlated the perpendicular and the in-plain strain with respect to each composition, it is concluded that the elastic constants of the studied ternary alloys follow a linear trend between those of GaN and InN binary compounds. These results have been confronted to and supported by measurements of the strain state in low and very high In concentration InGaN/GaN heterosystems by different research teams. Moreover, we propose a model of the critical thicknesses, which delimitates a transition state between pure pseudomorphic and completely relaxed layers as the In concentration on InGaN/GaN varies in the complete range of compositions (0<x<1). Our observations agree with the linear relationship between lattice spacing and the alloy composition demonstrated for free-standing InGaN (i. e. nanowires [2] or non-coalesced island growth) and denies the idea that this relation is subjected to a noticeable bowing parameter. 1. 2.
V. Lebedev, V. M. Polyakov, S. Hauguth-Frank, V. Cimalla, Ch. Y. Wang, G. Ecke, F. Schwierz, A. Schober, J. G. Lozano, F. M. Morales, D. González and D. González, J. Appl. Phys. 103, in press (2008). T. Kuykendall, P. Ulrich, S. Aloni and P. Yang, Nature Materials 6, 951 (2007)
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Convergence of microscopy techniques for nanoscale structural characterization: an illustration with the study of AlInN A. Mouti1, S. Hasanovic1, M. Cantoni1, E. Feltin2, N. Grandjean2, P. Stadelmann1 1. Interdisciplinary center for electron microscopy (CIME), EPFL, Lausanne, Switzerland 2. Laboratory of Advanced Semiconductors for Photonics and Electronics (LASPE), EPFL
[email protected] Keywords: III nitrides, TEM, SEM, AFM, Alloy, Dislocation, InAlN, AlInN, GaN
AlInN is a wide band gap group III nitride alloy, it is gaining interest nowadays because of its ability to be lattice matched to GaN, thus reducing defect creation. Recent research [1] has allowed the growth of good quality AlInN /GaN heterostructures on sapphire. However, AlInN is quite a complex alloy to grow and characterize, partly because of the high difference in ionic radii between Aluminum and Indium. The high dislocation density inherited from the lattice mismatch between GaN and Sapphire adds considerable complexity. Therefore, many physical phenomena can occur in this alloy, making it an interesting material for the microscopist. We present a broad structural study of AlInN to characterize its order and disorder, defects, interfaces, e-beam sensitivity. Weak beam dark field has been used to characterize dislocations and has been coupled with AFM to study surface pits “Figure 1”. EDX microanalysis in TEM was used to quantify compositional fluctuations with a 0.5% accuracy and a 3 nm spatial resolution “Figure 2”. HRTEM and CBED and precession diffraction coupled with JEMS simulations [2] are used to probe atomic structure order as well as to describe stress distribution. High resolution SEM was used to describe layer surfaces on a large 2 dimensional scale, providing information on thickness variations. When putting together the pieces of the puzzle brought by each technique, we show that we obtain quite a clear image of the alloy, thus providing valuable elements to the understanding of its optical properties. Results include and are not limited to stress induced composition variation around dislocations in agreement with theory, spiral growth composition fluctuation “Figure 2”, accurate description of pit dimensions and a better understanding of their origins “Figure 1”, interface disorder brought by dislocations, spiral growth induced thickness variation and an accurate description of defect creation and structural degradation with the increase of layer thickness. 1. 2.
J.-F. Carlin and M. Ilegems, Appl. Phys. Lett. 83, 668 (2003) P. A. Stadelmann, Ultramicroscopy Volume 21, Issue 2, 1987, Pages 131-145.
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Figure 1: Surface pits originating from screw component threading dislocations shown by weak beam dark-field thickness fringes
Figure 2: EDX Mapping showing periodic Indium peaks around a threading dislocation
Figure 3: HRTEM image of a mixed threading dislocation crossing a GaN/AlInN interface, along with the projected shear stress it is causing calculated by geometric phase analysis
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TEM analyses of microstructure and composition of AlxGa1-xN/GaN distributed Bragg reflectors A. Pretorius, A. Rosenauer, T. Aschenbrenner, H. Dartsch, S. Figge, and D. Hommel Institute of Solid State Physics, University of Bremen, Otto-Hahn-Allee 1, 28359 Bremen, Germany
[email protected] Keywords: AlGaN/GaN DBRs, TEM, STEM, STEM-simulation
In this work we analyse the microstructure as well as layer thicknesses and Al concentration of AlxGa1-xN/GaN distributed Bragg reflectors (DBRs) by transmission electron microscopy. To gain a high reflectivity of the DBR, a high refractive index contrast Δn is necessary, i. e. high Al content x. On the other hand, this leads to built-in stress. Due to this, cracks and/or dislocations can form and thus the reflectivity can be reduced. An optimum stress compensation and the adjustment of the wavelength reflected by the DBR needs a precise control of the layer thicknesses and the Al content. The analysed DBR was grown by metal organic vapour phase epitaxy (MOVPE). About 1 µm AlGaN of an intended concentration x = 0.25 was deposited on top of (0001) oriented Al2O3. This was followed by 40 mirror pairs of GaN and AlGaN with intended concentration x = 0.45 and nominal thicknesses of 50 nm. Finally, a GaN cap layer was deposited. The concentrations of the AlGaN buffer and DBR layers were chosen in order to obtain an overall stress compensated structure. Cross-sectional TEM samples were prepared and analysed in a TITAN 80/300 equipped with an imaging CScorrector and operated at 300 kV. Optical microscopy and TEM exhibited very few cracks in the DBR (Figure 1). Scanning electron microscopy (SEM) analysis revealed that no cracks penetrated to the surface, i. e. all are overgrown. In addition, few voids were observed by TEM analysis inside the DBR. Evaluation of the layer thicknesses from STEM images reveals a slowly varying, approx. constant AlGaN layer thickness (Figure 2). In contrast, the GaN layer thickness varies significantly. Using high resolution STEM (HRSTEM, Figure 3) it was found that the GaN-AlGaN interfaces are comparably sharp (interface width ~2.2 nm), whereas the AlGaN-GaN interfaces are blurred (interface width ~6.5 nm). The Al concentration x of the DBR layers was also obtained from Z-contrast images. The image intensity I was simulated in dependence of x and TEM specimen thickness t using the STEMsim program [1]. For t exceeding 20 nm, the ratio IGaN/IAlGaN was found to be nearly independent of the actual value of t, i. e. only to depend on x. On basis of the simulated intensity ratio, the Al concentration was extracted from the measured Zcontrast images and is displayed in Figure 4. The derived concentrations of ~0.43 for the AlGaN DBR layers and ~0.21 for the AlGaN buffer layer are in very good agreement with concentration values obtained from X-ray diffraction (0.45 and 0.21, respectively).
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In summary, the investigated AlGaN/GaN DBR is of good structural quality. The Al concentrations of the buffer and DBR layers are very close to the intended values. Nevertheless, the GaN layer thicknesses vary considerably, which may reduce the reflectivity of the DBR. 1. 2.
A. Rosenauer, M. Schowalter, Springer Proceedings in Physics: Microscopy of Semiconducting Materials (MSM) conference 2007, Cambridge. Financial support by the Deutsche Forschungsgemeinschaft (DFG Grant no. FOR 506) is gratefully acknowledged.
Figure 1. Z-contrast image of an overgrown crack.
Figure 3. HRSTEM image showing a sharp interface between GaN-Al0.45Ga0.55N, but a blurred interface between Al0.45Ga0.55N-GaN.
Figure 2. DBR layer thicknesses as obtained from Z-contrast images.
Figure 4. Concentration profile of the DBR obtained from a Zcontrast image.
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TEM study of quaternary InGaAsN/GaAs quantum well structures grown by molecular beam epitaxy T. Remmele1, M. Albrecht1, I. Häusler1, L. Geelhaar2, H. Riechert2, H. Abu-Farsakh3, and J. Neugebauer3 1. Institut für Kristallzüchtung, Max-Born-Str. 2, D-12489 Berlin, Germany 2. Qimonda (formerly Infineon Technologies Corporate Research Photonics), Otto-Hahn-Ring 6, D-81730 München, Germany 3. Max-Planck-Institut für Eisenforschung GmbH, Max-Planck-Straße 1, D-40237 Düsseldorf, Germany
[email protected] Keywords: InGaAsN, composition analysis, growth kinetics
In this study we report on the compositional analysis of quaternary InGaAsN/GaAs quantum well (QW) structures based on transmission electron microscope (TEM) investigations accompanied by theoretical ab-initio calculations. The pronounced band gap bowing by adding nitrogen to GaAs allows to reach the telecommunication wavelength range around 1.3–1.55 µm while staying lattice matched to the GaAs substrate by incorporating In. A sample series of InxGa1-xAs1-yNy with different nominal concentrations of In and N in the range of x = 0.1 – 0.4 and y = 0.01 – 0.05, respectively, were grown by radio-frequency plasma assisted molecular beam epitaxy and has been investigated by TEM. We used strain analysis of high-resolution TEM micrographs in combination with dark field imaging of the chemical sensitive (002)-reflection to evaluate the composition in the QW [1]. The evaluation of the dark field intensity depends on the scattering factors. For alloys the scattering factors are generally derived from atomic scattering factors using Vegard‘s rule. In recent publications [2,3] different scattering factors for InGaAs and GaAsN has been published. In these works the static atomic displacements [4] around individual In or N atoms, computed according to density functional theory, has been taken into account for the estimation of the scattering factors. However, in our evaluations we get a reasonable well correspondence between nominal and experimental values using scattering factors derived by Vegard‘s rule. The reason of this discrepancy is currently under investigation and will be reported. Figure 1 shows a dark field image of a GaAsN, an InGaAsN, and an InGaAs QW grown under identical conditions in one growth run. As shown in figure 2 we find at the bottom and also at the top interface of the InGaAsN QW a nitrogen enriched layer where the In concentration has its rising or trailing edge. The N-enriched layers occur in QW with an In concentration above x=0.2. The presence of these layers were also observed semi-quantitatively in energy filtered TEM measurements. On the ternary GaAsN QW no such enriched layers are present. The formation of the N-enriched layers is controlled by the surface kinetics during the growth. Ab-initio calculations show that while in bulk samples In-N pairs are
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energetically preferable due to strain reduction, at the surface, where strain relaxation occurs, the chemical environment dominates the energetic relations. The result is a repulsive In-N interaction and a high kinetic barrier for N-diffusion to subsurface positions at the growth surface which lead to the observed In-N anti-correlation in InGaAsN quantum wells. 1. 2. 3. 4.
V. Grillo, M. Albrecht, T. Remmele, H. P. Strunk, A. Y. Egorov, and H. Riechert, Journal of Applied Physics 90 (2001), p. 3792. A. Rosenauer, M. Schowalter, F. Glas, and D. Lamoen. Physical Review B (Condensed Matter and Materials Physics), 72 (2005), p. 85326. K. Volz, O. Rubel, T. Torunski, S. D. Baranovskii, and W. Stolz, Appl. Phys. Lett. 88 (2006), p 81910. F. Glas, Physical Review B, 51 (1995), p. 825.
Figure 1. Dark field image with (002)-reflection of three quantum wells with different compositions (InGaAs, InGaAsN, GaAsN). The nominal concentrations of In and N in the quantum wells are x=0.39 and y=0.02, respectively.
Figure 2. Evaluated N-concentration (left) and In-concentration (right) of the quantum wells in figure 1. N-riche interface layers are observed for the quaternary InGaAsN QW.
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Lattice polarity and capping of GaN dots studied by Z-contrast imaging J.L. Rouviere1, C. Bougerol2, J. Coraux2, B. Amstatt2 , E. Bellet-Amalric2 and B. Daudin2 CEA-Grenoble 17 rue des Martyrs 38054 Grenoble FRANCE 1. INAC/SP2M/LEMMA 2. CEA-CNRS group "Nanophysique et Semiconducteurs"
[email protected] Keywords: nitride, quantum dot, polarity, STEM
In this work we show the interest of using Z-contrast images for analysing GaN quantum dots. Firstly it is shown that the polarity of GaN quantum dots can be determined by using a probe Cs-corrector. Secondly, the way GaN quantum dots are capped with AlN layers is analyzed. Samples were grown by Molecular Beam Epitaxy (MEB). GaN quantum dots were deposited on or embedded in AlN layers [1]. For the capping study, different thickness of AlN capping layers where deposited on GaN dots (Figure 2). Cross-sections were realised by mechanical polishing and ion milling. A FEI-TITAN microscope equipped with a CEOS probe Cs-corrector and working at 300kV was used for the Z-contrast images. The polarity of two kinds of GaN quantum dots has been determined : GaN dots grown on (0001) AlN surface (Figure 1) and GaN dots grown on (1100) AlN (not shown here). (0001) GaN have a symmetric shape whereas (1100) GaN dots tend to have an asymmetric shape because in our Ga-rich growh condition, growth is faster along [0001] than along [0001] [2]. Figure 1 outlines how the Ga-polarity has been determined, that is to say how the GaN bonds parallel to [0001] (or GaN bond vectors) are oriented with respect to the surface normal of the layers. In Figure 1f, N atomic columns are not visible as clear white spots, because the contrast of N columns (Z=7) is broad and faint compared to Ga column (Z=31). However the location of N columns can be determined by comparing the regions above and below the central Ga atoms, which is equivalent at looking the surroundings of positions N2 and N4. Tunnels positions are clearly determined. In Figure 2, Z-contrast images determine at once the shape of the AlN capping layer. For AlN capping layer having an equivalent thickness smaller than 4 monolayers (ML), AlN wet the GaN dots. For AlN thicknesses between 4 ML and 14 ML, AlN fills the gap between GaN layers. Above 14 ML, the surface of the AlN capping layer is again flat [1]. 1 2
J Coraux, B Amstatt, JA Budagoski, E Bellet-Amalric, JL Rouvière, V Favre-Nicolin, MG Proietti, H Renevier, B Daudin Phys. Rev. B 74 (2006) p. 195302 J.L. Rouvière, C. Bougerol, E. Bellet-Amalric, B. Amstatt, B. Daudin, submitted to APL
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Figure 1 (a) High Resolution STEM image of a GaN at the surface of an AlN layer. (b) Atomic model of the wurtzite GaN projected on the (2,-1,-1,0) plane. In projection, two rectangular unit cells can be chosen. The left one is chosen in the other figures. (c) The left unit cell of a GaN crystal having a Ga-polarity orientation (surface normal is supposed to point upwards). (d) The left unit cell with a GaN crystal have a N-polarity orientation. (e) Image extracted from figure 1a). (f) Average of 10 images similar to figure 1e. N atomic column positions N1 and N2 (respectively N3 and N4) of a Gapolar crystal (respectively N-polar crystal) have been added to the picture. (g) Same image a figure 1f but with a slightly different contrast. The atomic structure of a Gapolar crystal has been added. Tunnels are clearly located. The difference between the regions above and below the central Ga atom of the rectangular cell is also clearly visible : the region below is darker than the region above, so the N-atomic column is situated above the central Ga atom.
Figure 2 Four Z-contrast images showing how AlN covers GaN dots for four different amounts of AlN : GaN material appears brighter than AlN. The numbers on the left of the picture give the number of equivalent AlN monolayers (ML) deposited on GaN dots. Above the AlN capping layer a thin GaN layer (2 ML) has been deposited in order to determine the shape of the AlN capping layer. On the top picture (4 ML) AlN wets GaN dots whereas on the bottom picture (14ML) AlN fills the gaps between GaN dots.
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Investigation of swift ions damage in wide band gap wurtzite semiconductors S. Mansouri, I. Monnet, H. Lebius, G. Nouet and P. Ruterana CIMAP, UMR 6252 CNRS-ENSICAEN-CEA- UCBN, 6, Boulevard du Maréchal Juin, 14050 Caen Cedex
[email protected] Keywords: nitride semiconductors, AlN, GaN, InN, swift ions irradiation, ion tracks, band gap smearing, optical properties, TEM, nanodots, AFM
The wide band gap wurtzite semiconductors (AlN, GaN, InN) constitute an interesting family of materials with an extensive range of micro-electronic and optoelectronic applications. Moreover, they are close to ceramics, therefore fabricated devices are also due to operate in harsh environments (space, underwater, inside motors or nuclear reactors, …). As the data on the effect of irradiation by swift heavy ions is still scarce (GaN, AlN) or completely inexistent (InN), their behaviour in such an environment is not known or is still controversial. The case of AlN is particular since it also has been investigated in the form of polycrystalline material, starting at the time when it was only considered for the fabrication of ceramic based devices. In good quality monocrystalline material, only AlN, and GaN have started to be investigated. Indeed, most of the work has been carried out under ion bombardment at medium range energies (~100 keV) for which the contribution of the electronic energy loss processes is rather negligible in the formation of defects. In the high energy range, wurtzite GaN layers were irradiated at 300K with 200 MeV 197Au16+ (Se = 34 keV/nm) and tracks were identified by TEM with a diameter of about 10nm. They were characterized by a structural disordering but no amorphous phase formation was reported [1]. More recently, damage accumulation was analyzed with 230 MeV Pb27+ at room temperature and at 513K. From RBS spectra, it was deduced that rapid damage accumulation and efficient erosion occur at a rather low value of the electronic energy deposition [2]. We can also quote irradiations with 150 MeV Ag12+ [3] and 100 MeV Ag8+ and O7+[4]. Probably due to the lack of good quality material, InN has not yet been studied and no report on the material behaviour under high energy ions is yet available. Therefore, it is of fundamental interest to carry out a systematic analysis of the electronic excitation in the energy deposition processes into these high expectation materials. Our investigations carried out with Xe (22.5 keV/nm) under grazing incidence show the formation of arrays of nanodots by single ions (Figure 1a). Moreover, a number of experiments have been also carried out with Pb ion beam on TEM thin foils, the first results indicate that there is a difference between AlN, GaN. No tracks were observed in AlN, whereas nanometer crystalline tracks are formed in GaN (Figure 1b).
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The evolution of the optical properties versus the ion fluence has also been investigated, for GaN, it has been found that, as the ion dose is increased above 10exp13 at*cm-2, the band gap is rapidly smeared out due to defects accumulation.
Figure 1 : a) arrays of nanodots in GaN irradiated at 1° xith Xe 0,7Me/u; b) Small tracks in GaN 0,5 MeV/u Pb 1. 2. 3. 4.
S.O. Kucheyev, H. Timmers, J. Zou, J.S. Williams, J. Jacadish and G. Li, J. Appl. Phys. 95, (2004)5360 C.H. Zang, Y. Song, Y.M. Sun, H. Chen, Y.T. Yang, L.H. Zhou and Y.F. Jin, Nucl. Instr. and Meth.B, 256, (2007)199 N. Sathish, Nucl. Instr. and Meth.B, 256, (2007)281 V. Suresh Kumar, Nucl. Instr. and Meth.B, 244, (2007)145
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A TEM analysis of the damage formation in thin GaN and AlN layers during rare earth ion implantation at medium range energy F. Gloux, M.P Chauvat and P. Ruterana CIMAP, UMR 6252, CNRS-ENSICAEN-CEA-UCBN, 6, Boulevard Maréchal Juin, 14050 Caen, France
[email protected] Keywords: GaN, ion implantation, nanocrystallisation, basal stacking faults, prismatic stacking faults, rare earth doping, TEM
Rare earth (RE) doped GaN has been the subject of research over the past few years due to promising light emission applications [1]. In optoelectronics, particularly, interesting results have been recently obtained, for instance laser action at room temperature (RT) have been reported in Eu doped GaN [2]. These applications were achieved using GaN layers doped in situ by molecular beam epitaxy. Ion implantation as a doping technique has interesting advantages, i.e. control of the doped area and easy lateral integration. However it causes a large amount of structural defects in the implanted layers. As the GaN fusion temperature is very high, post-implant annealing performed for RE activation should be carried out well above 1000°C. During the last few years, a renewed interest has been shown in understanding the structural damage build-up in the implanted GaN by a number of recent reports [4-5]. Recently, an improved optical activation of REs was obtained for annealing temperatures as high as 1300°C [3] using an AlN capping. It has been shown that the 'amorphous layer' that usually forms inside the implanted layer at high fluences was reported to be instead made of broken crystals [6]. These results demonstrate that the knowledge of the implantation damage build-up in GaN has not yet been completed. In this work, the nature of the damage in GaN implanted by rare earth ions at 300keV and room temperature has been investigated by transmission electron microscopy versus the fluence, from 7x1013 to 2x1016at/cm2, using Er, Eu or Tm ions implantation. The density of point defect clusters has been found to increase with the fluence. In GaN, from about 3x1015at/cm2, a highly disordered 'nanocrystalline layer' (NL) starts to form at the GaN surface. Its structure exhibits a mixture of voids and misoriented nanocrystallites (Figure 1). Basal stacking faults (BSFs) of I1, E and I2 types have been observed from the lowest fluence, they are I1 in majority. Their density increases and saturates when the NL appears. Many prismatic stacking faults (PSFs) with Drum atomic configuration have been observed [7, 8]. The I1 BSFs are shown to propagate easily through GaN by folding from basal to prismatic planes thanks to the PSFs (Figure 2: position of the PSF indicated by arrow). When implanting through a 10nm AlN cap, the NL formation threshold goes up to about 3x1016at/cm2, which support the point that the AlN cap plays a protective role against the damage formation
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in GaN up to the highest fluences [3]. The flat surface after implantation and the absence of SFs in the AlN cap indicate also the high resistance to the damage formation. 1. 2. 3. 4. 5. 6. 7. 8.
D. S. Lee and A. J. Steckl, Appl. Phys. Lett. 80, (2002)1888 J. H. Park and A.J. Steckl, Appl. Phys. Lett. 85, (20044588 Lorenz, U. Wahl, E. Alves, S. Dalmasso, R. W. Martin, K. P. O’Donnell, S. Ruffenach and O. Briot, Appl. Phys. Lett. 85, (2004)2712 C. Liu, B. Mensching, M. Zitler, K. Volz and B. Rauschenbach, Phys. Rev. B 57, (1998)2530 S. O. Kucheyev, J. S. Williams, C. Jagadish, J. Zou and G. Li, Phys. Rev. B 62, (2000)7510 F-R. Ding, W-H. He, A. Vantomme, Q. Zhao, B. Pipeleers, K. Jacobs, I. Moerman , Materials Science in Semiconductor Processing 5, (2003) 511 P. Vermaut, P. Ruterana and G. Nouet, Appl. Phys. Lett. 74, (1999) 694 F. Gloux, T. Wojtowicz, P. Ruterana, K. Lorenz, and E. Alves, J. Appl. Phys. 100, (2006)073520
Figure 1: a)Nanocrystalline layer with voids (arrow) in the surface area of a GaN subsequent to 4*e1015 at/cm2 Eu implantation at 300 KV. b)The corresponding TFF
1 nm Figure 2: An I1 basal stacking fault which folds to and from the prismatic plane, the positions of the stacking fault segments have been marked, see the cubic stackings.
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Characterization and modelling of semiconductor quantum nanostructures grown by droplet epitaxy D.L. Sales1, J.C. Hernandez2, P.A. Midgley2, A.M. Beltran1, A.M. Sanchez1, T. Ben1, P. Alonso-González3, Y. Gonzalez3, L. Gonzalez3 and S.I. Molina1 1. Departamento de Ciencia de los Materiales e I. M. y Q. I., Universidad de Cadiz, Campus Univ. Rio San Pedro, 11510 Puerto Real, Spain 2. Department of Materials Science and Metallurgy, University of Cambridge, Pembroke Street, Cambridge, CB2 3QZ, UK 3. Instituto de Microelectronica de Madrid (IMM-CNM, CSIC), Isaac Newton, 28760 Tres Cantos, Spain
[email protected] Keywords: quantum dots, quantum rings, electron tomography.
One of the main requirements for the development of quantum computing is the ability of manufacturing single photon emitters. Self assembled semiconductor nanostructures are ones of the most reliable candidates for this application, once control of size and spatial location is achieved. The droplet epitaxy technique allows growth of low density distributions of nanostructures with good optical quality, constituting a promising method for single photon devices and quantum applications [1]. In this work two samples have been analysed: (i) InAs quantum dots (QDs) nucleated on GaAs nanoholes that were formed by the crystallization of Ga droplets under As atmosphere [1], and (ii) GaSb quantum rings (QRs) formed by the crystallization of Ga droplets under Sb atmosphere. In this case, surface QR were also grown after growth of a 20 nm thick GaAs cap layer. In both cases, after the capping growth process, it was observed the formation of surface GaAs mounds that, as shown in this work, would be related to the location of the buried nanostructures. The capacity of transmission electron microscopy (TEM) to extract fine-scale information about the internal structure has been essential to understand the growth of these devices. Conventional TEM allows the characterization of structural defects which can appear during the capping of the nanostructures reducing the photoluminescence emission. It is observed in CTEM images that every buried mound is correlated with a mound on surface (see Figure 1). This fact enables the density of buried nanostructures to measured, directly related to the number of emitters, by analysing the sample surface. Three-dimensional scanning transmission electron microscopy (STEM) has been used to characterize GaSb QRs grown on a GaAs mounds. Reconstructed STEM tomograms show contrast arising from structural defects. These defects (probably dislocations) run from the buried mound to the surface mound. The 3D reconstruction (Figure 2) allows a detailed characterization of the structure, including the shape and size of the QR structure and nanohole, and the spatial shift between the buried and the superficial QR.
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Additionally, the strain state of the structures has been calculated from elasticity theory using finite element analysis. The places where the nanostructures nucleate correspond with those with larger in-plane lattice parameter, i.e. lower elastic energy. 1. 2.
P. Alonso-Gonzalez, B. Alen, D. Fuster, Y. Gonzalez and L. Gonzalez, Appl. Phys. Lett. 91 (2007) 163104. This work has been supported by MEC (TEC2005-05781-C03-01 and 02/MIC), SANDiE and ESTEEM European networks (Contracts NMP4-CT-2004-500101 and 026019 ESTEEM), and Junta de Andalucía (PAI research group TEP-120).
Figure 1. CTEM image of a sample where InAs QDs have been grown on GaAs mounds. Two buried mounds are shown correlated with ones on the surface.
Figure 2. Surface render of a STEM tomographic reconstruction of a GaSb QR grown on a GaAs mound to reveal the complex shape of the QR and a nanohole.
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Transmission Electron Microscopy Investigation of Self-Organized InN Nano-columns H. Schuhmann1, C. Denker1, T. Niermann2, J. Malindretos1, A. Rizzi1, M. Seibt1 1. IV. Physikalisches Institut, Georg-August-Universität Göttingen, D-37077 Göttingen, Germany 2. now at: Institut für Optik und Atomare Physik, Technische Universität Berlin, D-10623 Berlin, Germany
[email protected] Keywords: InN, nano-columns, wires
Semiconductor InN nano-columns have electronic properties which make them a promising candidate for novel photovoltaic devices. Molecular-Beam Epitaxy (MBE) grown InN on p-Si(111) substrates forms selforganized InN nano-columns under suitable conditions [1]. In a bimodal growth regime nano-columns with lengths of up to 2 microns and diameters down to 20 nm were observed. Because the shape of the nano-columns sensitively depends on growth parameters their crystal quality and characteristics were studied by Transmission Electron Microscopy (TEM). High-resolution TEM of a cross-section specimen at the interface between the crystalline InN nano-columns with a growth direction along the [0001] direction and the silicon substrate indicates the existence of an amorphous surface layer with a thickness of about 2nm. Additionally, a predicted tilt [2] between the nano-columns and the silicon substrate of about 23 degrees was observed. In order to study the crystal quality of the InN nano-columns along the rod axis projection, Focused Ion Beam (FIB) was used to prepare a TEM sample from an individual, specially selected long nanocolumn obtained from the bimodal growth regime. High-Resolution TEM micrograph (see Figure 1) of this slice showed a high crystal quality with no defects or oxidation. In particluar, no sub-surface microstructure or chemistry has been observed indicating homogeneous InN nano-columns. In order to analyse a possible distribution of axial rotations of the columns against the silicon substrate it is necessary to study a large ensemble of nano-columns along the growth direction. This is carried out by a bevel-view specimen preparation. Here a TEM lamella was cut out by an angle of 10° to the interface by the FIB (see Figure 2). As is seen in Figure 3 it is additionally possible to study the development of the number and diameter of the nano-columns with the distance to the silicon interface. 1. 2.
Christian Denker, Joerg Mailindretos, Florian Werner, Friederich Limbach, Henning Schuhmann, Michael Seibt, Angela Rizzi, Self organized growth of InN-nanocloumns on pSi(111) by MBE, physica status solidi, accepted (2008) Christian Denker, Diploma thesis, Universität Göttingen (2007)
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Figure 1. HRTEM micrograph of a single InN nano-column prepared by focussed ion-beam technique. The crosssection is parallel to the [0001] growth direction of the nanocolumns.
Figure 2. Sketch of a bevel-view sample preparation using focused ion beam
Figure 3. TEM micrograph of the bevel-view specimen. The silicon substrate can be seen in the lower left corner.
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Investigations on a dilute magnetic semicondutor (Ga1-xMnxAs) by conventional TEM and EELS M. Soda1, U. Wurstbauer, M. Hirmer, W. Wegscheider, J. Zweck 1. University of Regensburg, Institute of Experimental and Applied Physics, Universitätstrasse 31, 93040 Regensburg, Germany
[email protected] Keywords: GaMnAs, EELS
The Dilute Magnetic Semiconductor (DMS) Ga1-xMnxAs offers good prospects for the integration of ferromagnetic and semiconducting properties for use in future “spintronic” applications [1]. GaMnAs is grown under low temperature conditions, leading to the incorporation of various defects. The location and type of these defects determines the Curie temperature (Tc). Low temperature annealing generally increase Tc, while annealing at higher temperatures leads to a decrease of Tc [2, 3]. Therefore it is important to investigate the structure of GaMnAs in various states of annealing. The investigated samples consists of 20-80 nm thick Ga1-xMnxAs layer grown by Molecular Beam Epitaxy (MBE) with x varying between 1% and 6%. As grown and annealed materials were compared. From 002 DF images (see Figure 1) a homogeneous doping with Mn can be proven. In the annealed one, as expected, an about 2 nm thick Mn-oxide layer at the surface is present (see Figure 2). Its composition was investigated by EELS, clearly indicating an oxygen and a Mn peak. It is known that the temperature treatment leads to a diffusion of Mn to the surface [4]. The composition of the thin dark strip (see arrow in Figure 2) on the top of the GaMnAs layer remain unidentified. An excessive annealing (68h @ 230°C) on a sample was performed and the results are shown in Figure 3. In addition to the Mn-oxide layer and the dark strip, clusters are formed in about 8 nm deep cavities. So far it has not been possible to determine the composition or the lattice constant of the clusters due to their instability under the beam. Cooling the specimen at liquid nitrogen temperature is expected to help. For a concise interpretation of the results one has to bear in mind that during the ion milling the specimen suffers an increase in temperature. This increase may reasonably be assumed to be comparable to the annealing temperature (180- 250 °C), thus leading to a preparation artifact.
1. 2. 3. 4.
Y. Ohno et al, Nature 402 (1999), 790. S. J. Potashnik et al, App. Phy. Let. 79 (2001). J. Sadowski, Mat. Sci.-Poland 24 (2006). K. W. Edmonds et al, Phys. Rev. Let. 92 (2004), 037201
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a)
20 nm
b)
20 nm
Figure 1. 002 DF images show a homogeneous doped layer. a) as grown b) annealed.
a)
10 nm
Oxygen
Manganese
532 eV
640 eV
b)
5 nm
Figure 2. a) as grown sample, b) annealed sample with about 2 nm Mn-oxide layer, inset) EELS spectrum of the Mn-oxide layer.
20 nm
Figure 3. Micrograph of an excessive annealed Specimen. It can be seen a Mn-oxide layer of about 10 nm at the surface of the sample and cluster in 10 nm deep cavitys.
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About the determination of optical properties using fast electrons M. Stöger-Pollach University Service Center for Transmission Electron Microscopy, Technische Universität Wien, Wiedner Hauptstrasse 8-10, A-1040 Wien, Austria
[email protected] Keywords: optical properties, EELS, retardation
In the last years spatial and energy resolution of transmission electron microscopes (TEM) were improved tremendously by the invention of Cs-correctors and monochromators. Using electron energy loss spectrometry (EELS) in such TEMs makes the determination of optical properties of semiconductors accessible even for layers thinner than the wave length of light, where optical methods fail. It was long a time thought that simply using the Kramers-Kronig relations can give accurate information on the dielectric function and all the other optical constants. Frequently discrepancies between optical methods and EELS were found. The reason for this are retardation effects and Čerenkov radiation [1], emitted as soon as the speed of the probe electrons v is faster than the speed of light inside the sample c/n – with n as the refractive index: v>c/n. (1) If the problem of the zero loss peak (ZLP) removal is solved properly [2], there are a few possibilities to increase the accuracy of the determination of optical properties with fast electrons: (a) in the frame work of Kröger’s theory an iterative routine can be set up removing the relativistic effects off-line. (b) a reduction of the incident beam energy would change the unequal sign in Eq. (1) depending on the materials refractive index n. (c) an experimental difference method [3] can be set up subtracting the relativistic energy losses, which have an extremely narrow angular distribution. For the experimental demonstration we chose SiNx:H, because it shows slight relativistic effects for a 200 keV electron beam (v>c/n) but for 60 keV the speed of the electrons don’t exceed the speed of light inside the sample. Figure 1 demonstrates the ZLP removal problem for (left) an experiment with v
c/n on SiNx:H using several ZLP removal routines. It is found, that even if a correct band gap value is determined (by the power law fit method) the resulting data cannot be used for Kramers-Kronig Analysis (KKA). In Figure 2 we demonstrate the precision of the iterative routine removing relativistic effects by comparison with an experiment with v
E. Kröger, Zeitschrift für Physik 216 (1968) 115 – 135 M. Stöger-Pollach et al., Ultramicroscopy 107 (2007) 1178 – 1185 M. Stöger-Pollach, Micron doi:10.1016/j.micron.2008.01.23
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Figure 1: Comparison of ZLP removal routines for 60 keV (left) and 200 keV (right) experiments on 0.34 λ thick SiNx:H samples.
Figure 2: Experimental prove of the accuracy of an off-line correction based on Kröger’s theory. Left: Experiment performed with 60 keV avoiding retardation losses and right: experiment performed with 200 keV electrons and removing retardation losses off-line. Finally on both spectra KKA was performed in order to calculate the refractive indices of the two layers Figure 3: Calculation from 0-0.125 mrad showing that the surface intensities and the Cerenkov contributions can be excluded from the VEELS spectrum, if these small q-values can be excluded from the recording of the VEELS spectrum. Left: GaP using 60 keV, Right: GaP using 200 keV beam energy. B: Sketch of the principle experimental setup (From [2]). The Hochschuljubiläumsstiftung der Stadt Wien, contract Nbr. H-01585/2007, is kindly acknowledged for funding.
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Mapping uncompensated spins in exchange-biased systems by high resolution and quantitative magnetic force microscopy H. J. Hug, M. Marioni, S. Romer, I. Schmid and S. Romer Empa, Überlandstrasse 129, CH-8600 Dübendorf, Switzerland [email protected] Keywords: magnetic force microscopy, exchange bias effect, uncompensated spins
Magnetic Force Microscopy is an ideal tool to image magnetic stray fields emanating from surfaces but also from hidden interfaces of magnetic samples. A lateral resolution of 10nm is routinely obtained on flat samples. Tip calibration techniques were developed for a quantitative evaluation of the magnetic surface charge or surface dipole density from the measured MFM signal [1,2]. The origin of the exchange bias (EB) effect has been traced back to the existence of pinned uncompensated spins (UCS) in the antiferromagnet (AFM) or at its interface. However, the understanding of the underlying mechanism is still clouded by contradictory reports: Ohldag et al. [3] used XMCD to extract the direction and density of pinned UCS from the vertical shift of element specific magnetization loops. In their interpretation, the positive vertical shifts found for all investigated EB materials reflects a parallel coupling of the UCS to the spins of the ferromagnet (FM). Tsunoda et al. [4] used a similar experimental method but did not find any vertical shift in similar EB materials, although these materials showed a large EB-effect. Furthermore, Kappenberger et al. [5] (magnetic force microscopy) and Eimüller et al. [6] (XMCD PEEM) reported an antiferromagnetic coupling between the UCS and the FM spins. In order to clarify the role of the UCS for the EB-effect different perpendicular exchange-biased AFM/FM multilayers were prepared by magnetron sputtering. The FM layers consisted of CoPt and CoPd multilayers exhibiting a strong perpendicular magnetic anisotropy. For the AFM layers either CoO or MnIr was used. Magnetometry and magnetic force microscopy measurements were used advantageously to demonstrate the co-existence of pinned UCS that are parallel and antiparallel to the cooling field in metallic (IrMn) and oxidic (CoO) EB systems. We found that the exchange-bias-effect (EB-effect) is a result of pinned interfacial UCS, which are antiparallel to the spins of the ferromagnet. The often observed positive vertical shift of the magnetization loop after field cooling is due to pinned UCS that align parallel to the cooling field, but are of little importance for the EB-effect [7]. Given that most experimental techniques (including those used in our work [5,7]) measure the sum of both spin groups indiscriminately, contradictory results can arise. Our work may provide guidelines for the design of experiments that can correctly determine the densities of those UCS that do contribute to the EB effect. In further experiments, the distribution of density of the UCS was imaged on the length scales of single grains. A surprisingly strong fluctuation of the local UCS density
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(UCSD) was observed. A correlation between the UCSD and the local exchange field was performed. Clearly, a high UCSD results in a high local exchange bias field. Regions with an anti-biasing effect were found. Using grain-boundary engineering, exchange-biased materials without such regions could be fabricated that showed a substantially increased exchange bias effect. 1. 2. 3. 4. 5. 6. 7.
P. J. A. van Schendel, H. J. Hug, B. Stiefel, S. Martin, H.-J. Güntherodt, J. Appl. Phys. 88 (2000), p435 M. A. Marioni, N. Pilet, A.V. Asworth, R. C. O'Handley, H. J. Hug, Phys. Rev. Lett. 97 (2006), p027201 H. Ohldag et al., Phys. Rev. Lett., 87 (2001), p247201. M. Tsunoda et al., Appl. Phys. Lett. 89 (2006), 172501 P. Kappenberger, S. Martin, Y. Pellmont, H.J. Hug, J.B. Kortright, O. Hellwig, and E.E. Fullerton, Phys. Rev. Lett. 91 (2003), p267202 T. Eimueller et al., J. Appl. Phys. 85 (2004), p2310. I. Schmid, P. Kappenberger, O. Hellwig, M. Carrey, Eric E. Fullerton, H. J. Hug, Euro Phys. Lett. 81 (2006), p17001
Figure 1. MFM images of a CoO/CoPt-multilayer antiferromagnet/ferromagnet sample. Panels a) shows the domains of ferromagnetic CoPt-multilayer after zero-field cooling and after a field of 100mT (panel b) and 200mT (panel c) is applied. Saturation of the ferromagnetic layer is obtained at 300mT (panel d and e). Then, the pinned uncompensated spins at the antiferromagnet/ferromagnet interface become visible (panels d and e). The yellow lines visible in panels d) and e) are guides for the eye to indicate the walls of the ferromagnetic domain states after zero-field cooling and at 200mT.
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Ferroelectric materials and structures suitable for data storage: The role of microscopies in establishing preparation-microstructure-property relations D. Hesse1, M. Alexe1, K. Boldyreva1, H. Han1,2, W. Lee1, A. Lotnyk1, B.J. Rodriguez1, S. Senz1, I. Vrejoiu1, and N.D. Zakharov1 1. Max Planck Inst. of Microstructure Physics, Weinberg 2, D-06120 Halle, Germany 2. Dept. of Mater. Sci. and Eng., Pohang University of Science and Technology, Korea [email protected] Keywords: ferroelectrics, superlattices, nanostructures, microscopies
Among the various possible routes towards solid-state non-volatile random access memories (NV-RAMs) of high memory density, ferroelectric random access memories (FRAMs) offer very good prospects. Non-volatility is provided by the ionic character of the signal, and switching by an electric field is sufficiently fast and free from crosstalk. In view of these excellent prospects, ultrathin epitaxial films, multilayers, superlattices, and nanostructure arrays made from (anti-)ferroelectric perovskite or perovskitelike oxides as PbZrO3 (PZO), Pb(Zr,Ti)O3 (PZT) or (Bi,La)4Ti3O12 (BLT) are a most actual subject of research in materials science. Modification and tuning of interesting properties is possible by defect and orientation engineering in thin films, by variation of composition, layer thickness, strain or crystallographic orientation in multilayers and superlattices, and by choosing lateral size, pitch or crystallographic orientation of nanostructure arrays. Preparation by pulsed laser deposition (PLD) offers a number of advantages, especially in research. Ultrathin epitaxial films or strained superlattices (Figure 1), but also well-ordered, large-area arrays of ferroelectric nanostructures and nanocapacitors (Figure 2) can conveniently be prepared by PLD. The deposition of nanostructure and nanocapacitor arrays requires the application of appropriate nanoscale deposition masks made from metallic or well-ordered anodic alumina (AAO) layers. One of the main advantages of PLD is the possibility to extensively vary the deposition conditions in order to establish preparation-microstructure-property relations. Such relations can be studied in detail, comprising, e.g., the structure of 90° and 180° ferroelectric domain boundaries (e.g., [1]) and their role in switching processes (e.g., [2]). As will be shown, microscopies like AFM, SEM, piezoresponse-mode AFM (PFM), and (high-resolution) TEM are significant tools to study these relations. Plan-view and cross sectional (high-resolution) TEM images, as well as SEM and AFM topography images are used to characterize the microstructure of the substrates and the grown superlattices and nanocapacitor arrays (Figures 1,2), whereas PFM (Figure 2) and its recently developed switching-spectroscopy variant (SS-PFM) [2] are applied in studies of the local switching properties and domain dynamics. 1. 2.
C.-L. Jia, S.-B. Mi, K. Urban, I. Vrejoiu, M. Alexe, D. Hesse, Nat. Mat. 7 (2008) 57. S. Jesse, B.J. Rodriguez, S. Choudhury, A.P. Baddorf, I. Vrejoiu, D. Hesse, M. Alexe, M.A. Eliseev, A.N. Morozovska et al., Nat. Mat. 7 (2008) 209.
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Figure 1. (a,b) AFM topography images (3 x 3 μm2) of (a) a SrTiO3 (100) substrate with 0.1° miscut, and (b) a PZO/PZT(80/20) superlattice (overall thickness 100 nm; layer thickness 12.5 nm). (c,d) TEM cross section images of a coherently strained PZO/PZT(80/20) superlattice (overall thickness 150 nm; layer thickness 8 nm) on a PZT(20/80) buffer layer and a SrRuO3 electrode on a vicinal SrTiO3 substrate.
Figure 2. (a) SEM image of an as-prepared PZT nanoisland array with part of the ultrathin AAO mask. (b) TEM cross section image of a regular Pt/PZT/Pt nanocapacitor array. (c)-(f) AFM/PFM images (260 x 260 nm2) of a Pt/PZT/Pt nanocapacitor array: (c) topography; (d)-(f) piezoresponse, viz. (d) before switching, (e) after positive switching of two capacitors applying +3V, and (f) after negative switching of one of the two previously switched capacitors applying -3V.
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Electronic structures at Magnetic Tunnel Junction interfaces: EELS experiments and FEFF calculations K. March, D. Imhoff, G. Krill, C. Colliex Lab. Physique des Solides, UMR CNRS-Université Paris XI, 91405 Orsay, France [email protected] Keywords: Heterostructures, spintronic, EELS, multiple scattering calculation
Electron transfer between the electrodes in a Magnetic Tunnel Junction (MTJ) is spin-dependent and directly related to the relative orientation of the magnetization of the two ferromagnetic layers on each side of the insulator barrier. TMR (tunnel magnetoresistance) values of several hundreds of % have been measured for the present MTJ generation, however well below the theoretical expectations. Higher TMR values measured on the epitaxial Co(Fe)/MgO/Co(Fe) system, reflect the position of the Fermi level which lies in the d-band where the electrons with Δ5 symmetry are only partly spin-polarized. Bridging experimental results with theoretical modelling requires a deeper investigation of the hybridization with the d electrons of transition atoms at the interfaces. In-situ XPS measurements on the Fe/MgO [1] have revealed that the hybridization between the Fe(d) and O(p) electrons seems to be rather small. However, we have no information about the CoFe/MgO system except a recent work (CoFeB [2]). We have therefore addressed this issue by recording EELS spectra of the O-K edge at sub-nanoscale range and by simulating the fine structures with multiple scattering calculations (FEFF 8.4). Co0.6Fe0.4/MgO/Co0.6Fe0.4 epitaxial samples on MgO substrates have been grown by Molecular Beam Epitaxy (MBE) at LPM Nancy. The atomic structure of the interfaces has been checked by HRTEM observations on a FEG-TEM equipped with a Cs corrector at CEMES Toulouse. EELS experiments resolved both spatially (< 0.6 nm) and energetically (< 0.5 eV) have been performed with a STEM microscope at LPS-Paris-Sud Orsay (3D spectrum-image 128*128*1340 pixels or less, core-loss area study (elemental profiles, ELNES extraction, branching ratio,…). O K-edges, directly extracted from spectrum-images (Figure 1A), show that O2p and (Fe,Co)3d bonding (oxidized TM) as well as e-beam induced O2 desorption, manifest themselves by the appearance of pre-edge structures located between 528 and 535 eV. FEFF calculations of the O K-edge spectra in the last interface plane of MgO confirm the appearance of prepeak structures depending on the Co(or Fe) atom position (Figure 2). They also show that only a slight longitudinal shift of the bcc transition metal structure relative to the MgO one, may produce a noticeable modification of the preedge on a simple “perfect” TM-O interface. This is of prime importance to understand the origin of the high TMR reported in this system. Experimental O K-edges (Figure 1B extracted from “line” or “sum” spectra) exhibit very similar “shapes” at the “top” and “ bottom” interfaces, with weak features in the prepeak region more similar to the hollow than to the apex calculated profiles, in agreement with the published XPS data [2]. These first results will be further developed for improving our understanding of the electronic properties of MTJ interfaces and effects on transport properties.
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S. Yuasa et al, Appl. Phys. Lett. 89 (2006), p 042505. J.J. Cha et al., Appl. Phys. Letters 91 (2007), p 062516 N. Bonnet, N. Brun, C. Colliex, Ultramicroscopy 77 (1999), p 97 Thanks are due to S. Andrieu and colleagues at LPM Nancy for specimen elaboration, to E. Snoeck and V. Serin at CEMES, Toulouse for collaboration. K.M. acknowledges ALTIS for her CIFRE thesis grant. The EC programme ESTEEM is also acknowledged for its support.
Figure 1. Experimental EELS results for the O K-edge. (A) reference data (≤ 0.2eV/ch): (i) pure MgO substrate, (ii) irradiation damage in MgO (chronospectroscopy (50 x 2 s)), (iii) oxided FexCo(1-x). (B) (upper part) as-acquired line scan across the MgO/FeCo heterostructure (0.2 eV/ch) and (lower part) sum of 128 spectra along few decades of nm at different spatial beam positions (“bottom” and “top” interfaces, “center” of the barrier), (0.3 eV/ch).
Figure 2. (A) Different positions of Co(or Fe) atom are considered relative to the central O atom in the last plane of the MgO structure: in “apex” or in the “hollow” site position. (B) Corresponding theoretical O K-edges calculated using FEFF 8.4 (with 500 at.): (i), (ii) (“apex” position) and (iii) (“hollow” site). Spectra (ii) and (iii) were convoluted by a 1eV Gaussian (FWHM)) in order to take into account of experimental conditions.
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Stability and reaction of magnetic sensor materials studied by atom probe tomography G. Schmitz1, C. Ene2, H. Galinski1, V. Vovk1 1. Institut für Materialphysik, Westf. Wilhelms-Universität, Wilhelm-Klemm-Str. 10, 48149 Münster, Germany 2. Institut für Materialphysik, Universität Göttingen, Friedrich-Hundt-Platz 1, 37077 Göttingen, Germany [email protected] Keywords: GMR sensors, atom probe tomography, thermal reaction, Cahn-Hilliard-theory
State-of-the-art magnetic sensors usually consist of thin film multilayers with a periodicity of a few nanometers only. The significant contribution of interfaces and grain boundaries makes these nano-scaled materials inherently unstable and, even more puzzling, their high density of defects may accelerate atomic transport in an unpredictable way. Induced by the nanocrystalline microstructure of the sputtered thin films, reactions develop on a complex morphology which requires three-dimensional high resolution analysis as it is provided by atom probe tomography in a unique way. In the talk we will address the reaction of Cu/Py (Py stands for the soft magnetic Ni81Fe19 alloy) and of the Cu/Co model system. Compared to the frequently used Cu/Co sensors, multilayers of Cu/Py are distinguished by very low hysteresis which is of advantage for the design of position or orientation sensors. On the other hand, the thermal stability of these layer systems is rather low. What are the dominant structural mechanisms of thermal degeneration? We sputter-deposited thin film multilayers on top of tungsten substrate tips and studied the chemical structure after isochronal heat treatments in a temperature range from 150°C to 500°C [1]. Since the deposited layers reveal a columnar, nanocrystalline structure, one may expect that the dominant transport proceeds by grain boundary diffusion. Indeed, transport along grain boundaries is demonstrated by segregation zones in the tomographic volume reconstructions (see Figure 1a). In particular, grain boundaries inside Py appear enriched in Cu to surprisingly high amounts, up to 54at%. Such a Cu content would be way enough to insulate individual Py grains magnetically. However, these segregation zones develop only at temperatures higher than 350°C. In consequence, grain boundary transport cannot be made responsible for GMR degeneration starting already at much lower temperature. Due to the outstanding resolution of atom probe tomography, it is possible to detect even minor chemical modifications at the Cu/Py layer interface. So, we could observe in our experiments that the Cu concentration gradient at the interfaces decreases from about 1.5 nm-1 to about 1.0 nm-1 during 20 min annealing at 250°C due to a short-ranged volume diffusion. Such a modification seems to be negligible. However, in view of the typical single layer thickness of about 2 nm, it means that a considerable fraction of the
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volume has become already intermixed and in consequence the GMR amplitude has decreased significantly. Volume diffusion between insoluble metallic layers is quite surprising. In order to identify the physical basis of this process, we performed similar experiments with other model systems and found this short-ranged diffusion as a general behavior. This allows the conclusion that the mixing depth at the interface, although temperature dependent, must be controlled by a non-diffusive mechanism. Indeed, it is possible to describe the measured temperature dependence of the interface width by a Cahn-Hilliard ansatz (see Figure 1b), which predicts a maximum limit of the concentration gradient. So, as our main conclusion, we state that the thermal degeneration of the GMR amplitude is controlled by the concentration gradient term in the Gibbs energy. With atom probe tomography it is possible to prove this mechanism even on a length scale below one nanometer. C. Ene et al., Acta Mater. 5 (2005) 3383. P. Stender, C. B. Ene, H. Galinski, G. Schmitz, Intern. J. Mat. Res. (in press)
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Transmission Electron Microscopy Analysis of Tunnel Magneto Resistance Elements with Amorphous CoFeB Electrodes and MgO Barrier Michael Seibt1, Gerrit Eilers1, Marvin Walter1, Kai Ubben1, Karsten Thiel1, Volker Drewello2, Andy Thomas2, Günter Reiss2, and Markus Münzenberg1 1. Georg-August-Universität Göttingen, IV. Physikalisches Institut and Sonderforschungsbereich 602, D-37077 Göttingen, Germany 2. Fakultät für Physik, Universität Bielefeld [email protected] Keywords: MRAM, spin torque, amorphous/crystalline interface
Magnetic tunnel-junctions consisting of CoFeB/ MgO/ CoFeB trilayers have been of great interest in research just recently. Due to their high magnetoresistance they are a promising candidate for the fabrication of spintorque MRAM devices. For future writing concepts like current-induced magnetic switching, magnetic tunnel-junctions (MTJs) with thin barriers are necessary to provide sufficient high current densities. In such elements the tunnel magneto resistance (TMR) is strongly dependent on the electron transmission at the metal/oxide interfaces. Therefore the quality of the interfaces is of great significance and has to be optimized on the nanoscale. With the objective to correlate electrical transport properties (I-V characteristics, TMR) with the geometrical and chemical interface roughness, structural analysis was made by cross-sectional TEM and energy dispersive X-ray spectroscopy (EDX) in a Philips CM200-UT-FEG. Figure1 is a bright-field micrograph showing the multilayer stack grown on thermally oxidized silicon substrates. The heart of the device is the crystalline MgO barrier sandwiched between two CoFeB electrodes (see arrows in Figure 1). The atomic arrangement at the interfaces between electrodes and MgO barrier determine properties of spin transport across the barrier. The microstructure of this central part is characterized by a predominant interface roughness with a wavelength of about 10nm which may exceed the thickness of the barrier which is typically 2-2.5nm (Figure.2). Hence, for thicker TEM samples CoFeB electrodes and MgO barrier may overlap along the electron beam direction. Local crystallisation of the CoFeB electrodes is clearly revealed by high-resolution imaging. Due to the large roughness amplitude small-angle grain boundaries are observed in the MgO which predominantly grows in [100] direction. In addition, for non-working devices, cracking of MgO barriers is indicated by TEM. Although long-range ordering in CoFeB close to the MgO barrier is indicated in HREM images, the large interface roughness prevents a quantitative analysis on real devices. Films grown on crystalline MgO substrates with large atomically flat terraces indeed corroborate this hypothesis (see Figure 3). Post-growth thermal annealing indded shows that CoFeB cyrstallisation starts at interfaces to the MgO.
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Figure 1: Bright-field micrograph of a multilayer stack grown thermally oxidized silicon. Arrows indicate the functional part of the device, i.e. the crystalline MgO barrier between amorphous CoFeB electrodes.
Figure 2: Microstructure of the MgO barrier grown on a MnIr/CoFeB double layer as part of the multilayer stack shown in Fig.1. The predominant layer roughness has a wavelength of about 10nm and may exceed the barrier thickness of 2-2.5nm.
Figure 3: Amorphous CoFeB grown on (001) MgO substrate; the HREM image indicates long-range ordering in the CoFeB close to the substrate.
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Study of the intermixing of Fe–Pt multilayers by analytical and high-resolution transmission electron microscopy W. Sigle1, T. Kaiser1, D. Goll1, N.H. Goo1, V. Srot1, P.A. van Aken1, E. Detemple2, and W. Jäger3 1. Max Planck Institute for Metals Research, Heisenbergstraße 3, D-70569 Stuttgart, Germany 2. Universität des Saarlandes, Postfach 151150, D-66041 Saarbrücken, Germany 3. Christian-Albrechts-Universität zu Kiel, Kaiserstraße 2, D-24143 Kiel, Germany [email protected] Keywords: EDXS, interdiffusion, multilayers
Materials used for advanced magnetic data storage devices must offer high magnetic anisotropy, high saturation magnetization, and an easy magnetization direction which is perpendicular to the array plane. These demands are fulfilled by the L10-ordered FePt. The most promising preparation technique for thin L10-ordered FePt films is deposition of Fe–Pt multilayers on MgO followed by thermal annealing. In this study we performed analytical [1] and in-situ high-resolution transmission electron microscopy (TEM) [2] studies on these multilayers. Five layers of Fe (6.4 nm thickness) and five layers of Pt (5.6 nm thickness) were alternately deposited on MgO substrates by ion beam sputtering. Cross-sectional TEM specimens were prepared by ion-beam milling. For the analytical studies, the TEM specimens were annealed in a TEM at different temperatures between 360 and 735 °C using a heating stage. Chemical concentration profiles were obtained by energydispersive X-ray (EDX) analysis in a VG HB501UX dedicated scanning TEM equipped with a Noran SYSTEM SIX EDX. In-situ high-resolution TEM (HRTEM) was done in the JEOL ARM1250. Ordered FePt phases are identified by the 0.37 nm spacing of the ordered {100} planes in HRTEM and by (100) reflections in diffraction patterns. The fully L10-ordered FePt structure, as obtained after annealing at 735 °C, is shown in Figure 1. At intermediate annealing temperatures also L12 phases (Fe3Pt and FePt3) are detected (not shown). These phases can be distinguished from the L10 phase by the appearance of ordering along two orthogonal directions. Figure 2 shows three EDX linescans obtained from the as-prepared material and from material annealed at 565 °C and 735 °C. The drop of the amplitude of concentration modulations is readily visible. After final annealing no modulations are detected any more. A plot showing the modulation amplitude as a function of annealing temperature is presented in Figure 3. It is found that the modulation amplitude decreases within 3 steps near 400, 600, and 700 °C. Using literature data of diffusion of Fe and Pt in various FePt systems, we found that the three annealing steps are related to the onset of diffusion (i) within pure Fe and Pt, (ii) within L12 phases, and (iii) within the L10 phase.
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1. 2. 3.
T. Kaiser, W. Sigle, D. Goll, V. Srot, P. A. van Aken, E. Detemple, and W. Jäger, J. Appl. Phys. (2008) in print. D. Goll, A. Breitling, N. H. Goo, W. Sigle, M. Hirscher, and G. Schütz, J. Iron Steel Res. Int. 13, Suppl. 1 (2006) 97. We kindly acknowledge the help of Ute Salzberger for TEM specimen preparation and R. Höschen and K. Hahn for technical assistance.
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b Figure 1. High-resolution micrograph of the FePt multilayer after final annealing at 735 °C. The fringes belonging to ordered {100} planes are parallel to the MgO (100) substrate surface.
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Figure 3. Normalized amplitude of the Pt concentration modulation after annealing. The data show averages and standard deviations of the three central Fe/Pt layers. Dotted lines mark plateau regions
Figure 2. Pt (---) and Fe (—) concentration across the Fe–Pt multilayer stack. (a) as prepared and after annealing at (b) 565 °C and (c) 735 °C.
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Exploring structural dependence of magnetic properties in FePt nanoparticle by Cs-corrected HRTEM Z.L. Zhang1, J. Biskupek1, U. Kaiser1, U. Wiedwald2, L. Han2, P. Ziemann2 1. Materials Science Electron Microscopy, Ulm University, 89081 Ulm, Germany 2. Department of Solid State Physics, Ulm University, 89081 Ulm, Germany [email protected] Keywords: nanoparticles, Cs-corrector, HRTEM
Well-separated and size-controlled FePt nanoparticles are promising for magnetic data storages with ultra-high areal densities [1]. Chemically ordered L1o-FePt shows high magneto-crystalline anisotropy and an enhanced coercivity as compared to disordered FePt state. A measurement of coercivity (Hc) of FePt, showing strongly size dependence (shown in another paper), indicates a lower anisotropy when nanoparticle is small, which could be structural dependence. Newly-developed TEM techniques, Cscorrector and exit wave reconstruction, supplying less delocalized image, has enabled us to directly explore the structure of nanoparticle even with lateral dimension of only a few atoms[2]. In this work, an objective-lens Cs-corrected Titan FEI 80-300 microscope was used to explore the intrinsic structural characteristic of FePt particle with different sizes, and try to establish the structure-property relationship. In this study, all FePt particles were prepared by micellar approach [3-4], and followed by plasma etching treatment and reduction at 700°C. A series of representative HRTEM images with particle size ranges from 1.6 nm to 8.5 nm, recorded under a small Cs (less than 1 µm) and small over-focus, either cross-sectional (on MgO) or plan-view images, are shown in Figure 1. FFT from 2.3nm particle shows no superlattice reflections due to ordering. From the images, it is clearly seen that planar or line defects, i.e twin and dislocation, are dominant defects in these FePt particles, which lead to simply- and multiple-twinned particle. With the increase of particle size, the amounts of defect generally increase due to lattice mismatch, such as multiple twins in 7.2 nm and 8.5 nm particles, respectively. Epitaxial growth on MgO is possible only when the particle is smaller, i.e 1.6 nm, whereas it is lost as particle is larger, i.e. 6.0 nm, as a result, dislocations are usually introduced at interface or within the particle. The effective volume in the particle which contributes to magnetic moment could vary with the defect amount. Besides absence of ordering in smaller particles, the presence of these different types of defects in FePt particle may partially account for the observed magnetic properties. A statistic measurement and quantification about the relation of defect and particle size will be presented, further uncovering the intrinsic reasons for the demonstrated property. The stability of such FePt particle under electron beam and interface structure between FePt and MgO were investigated by both STEM and Cs-corrected HRTEM. Under STEM imaging, it is noted that Pt atom within a small FePt particle, i.e. 1~2 nm, is quite mobile. Closely examination on the behaviour of individual FePt under the
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beam reveals that atom cooperative movement in the particle occurs whereas the interface of FePt and MgO keeps quite stable due to a strong binding of Fe-O at the interface (Figure 2). Image simulations were conducted to prove the interfacial structure [5]. 1 2 3 4 5
O. Dmitrieva, et al, Phy.Rev B 76 (2007) 064414 R. M. Wang et al, Phy.Rev.Lett. 100, (2008)17205. U. Wiedwald et al , ChemPhysChem 6(2005)2522. G. Kästle et al, Adv.Funct.Mater., 13(2003)853. The authors thank the support from Germany Research Foundation DFG with SFB 569
Figure 1. HRTEM images of L1o-FePt nanoparticles recorded under Cs =1 µm , particle size is: (a) 1.6 nm, 2.3 nm with one twin and corresponding FFT is inserted, containing no superlattice reflections; (b) 4.0 nm, 2 twins; (c) 6.0nm, one edge dislocation in the particle; (d) 7.2 nm, five-fold twin, exit wave reconstructed phase image; (e) 8.5 nm, multiple twins. Note that the defect amounts are different in each particle.
Figure 2. A series of HRTEM images (FePt particle on MgO, viewed along [011], Wiener filtered) demonstrate under electron beam (0.042 A/mm2) the particle stability with exposure time, 0 s; 52 s and 280 s. It contains one twin boundary (indicated by arrow), atom cooperative movement along the boundary in the right part (three layer atoms) occurs.
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Understanding the Chemistry of Molecules in Nanotubes by Transmission Electron Microscopy A.N. Khlobystov1,2, M.W. Fay2,3 and P.D. Brown2,3 1. School of Chemistry, University of Nottingham, Nottingham NG7 2RD, UK 2. Nottingham Nanoscience and Nanotechnology Centre, Nottingham NG7 2RD, UK 3. School of Mechanical, Materials and Manufacturing Engineering, University of Nottingham, Nottingham NG7 2RD, UK [email protected] Keywords: nanotubes, fullerenes, TEM
Carbon Nanotubes (NTs) possess a spectrum of functional and exploitable properties such as high electrical and heat conductance, high tensile strength and superb thermal stability. One of the most fascinating physicochemical properties of nanotubes is their ability to host molecules due to their tubular topology [1-3]. The typical diameters of single-walled NTs range between 0.6 nm and 2 nm, and internal diameters of multiwalled NTs can reach 100 nm which makes carbon nanotubes ideal containers for a wide variety of molecular species. Several examples of templating action of nanotubes will be presented, and the underlying mechanisms of the nanotube-molecule interactions will be discussed. Most of the molecules interact with nanotubes via low-directional van der Waals forces. The most important parameter defining the efficiency of the NT-molecule interaction is the nanotube internal diameter. For spheroidal molecules, such as C60, the nanotube diameter controls the packing arrangement of the molecules (Figure 1). Most of the phases of C60 observed inside NTs do not exist in the free state and, therefore, are regarded as products of confinement in quasi-1D cavities of nanotubes. Various practical applications have recently been proposed for carbon nanotubes involved in non-covalent interactions with molecules, including sensors, photovoltaic devices and highly selective absorbents. One of the most fascinating developments in this area is the demonstration of chemical reactions inside nanotubes. The reactivity of molecules encapsulated in NTs can be controlled by this confinement. The nanotube in this case acts as a reaction vessel and a template for the reaction product, yielding unusual structures inaccessible by other means. One of the most illustrative examples is the controlled polymerisation of fullerene epoxide C60O recently carried out inside single-walled NTs [4,5]. Transmission electron microscopy (TEM) is the only direct method allowing visualisation of molecular structures inside carbon nanotubes. Positions and orientations of molecules can be visualised with atomic precision using TEM. The dynamic behaviour of molecules in nanotubes, such as translational and rotational motion, can also be monitored by taking time series TEM images. All this information is invaluable for the understanding of the effects of confinement in NT on the fundamental physicochemical properties of molecular species. The behaviour of the molecules may
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be altered significantly as a result of irradiation with e-beam in TEM. The extent of these secondary effects depends crucially on the accelerating voltage and the imaging conditions which have to be tuned carefully to avoid artefacts associated with this technique. 1. D.A. Britz and A.N. Khlobystov, Chem. Soc. Rev. 35 (2006), p. 367. 2. A.N. Khlobystov, D.A. Britz and G.A.D. Briggs, Acc. Chem. Res. 38 (2005), p. 901. 3. http://www.nottingham.ac.uk/nanocarbon/ 4. D.A. Britz, A.N. Khlobystov, K. Porfyrakis, A. Ardavan and G.A.D. Briggs, Chem. 5. 6.
Commun. (2005), p.37. A.N. Khlobystov in “Chemistry of Carbon nanotubes” American Scientific Publishers (2008), ISBN 1-58883-128-0 in press. We would like to acknowledge the EPSRC (UK), the European Science Foundation, the Royal Society and the University of Nottingham for financial support.
Figure 1. (a) Packing arrangement of fullerene C60 in the crystal. (b) High resolution transmission electron microscopy micrograph and a diagram of a C60@SWNT structure (nanotube diameter ~1.4nm). (c) Zigzag arrangement of fullerenes inside wide doublewalled carbon nanotube (inner diameter of the nanotube ~2.1nm).
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Electrical and mechanical property studies on individual low-dimensional inorganic nanostructures in HRTEM D. Golberg1,2, P.M.F.J. Costa1, M. Mitome1, Y. Bando2 and X.D. Bai3 1. Nanoscale Materials Center, National Institute for Materials Science (NIMS), Namiki 1-1, Tsukuba, Ibaraki, 3050044, Japan 2. World Premier International Center for Materials Nanoarchitectonics (MANA), National Institute for Materials Science (NIMS), Namiki 1-1, Tsukuba, Ibaraki, 3050044, Japan, 3. Institute of Physics, Chinese Academy of Sciences, Beijing, 100080, China [email protected] Keywords: STM-TEM, AFM-TEM, nanotubes, nanowires, nanobelts
The present contribution demonstrates the recent advances achieved by the Group in pioneering studies of electromechanical properties of diverse low-dimensional individual inorganic nanomaterials inside a transmission electron microscope. The measurements were performed in high-resolution field-emission transmission electron microscopes operating at 300 kV, namely JEOL-3000F and JEM-3100FEF, and on freestanding individual nanoscale objects by means of specialized piezo-driven STM-TEM and AFM-TEM holders commercialized by “Nanofactory Instruments AB”. A sharp etched Au tip, acting as the STM tip, or a Si cantilever, as the AFM probe, was assembled within the holders (Figure 1 and Figure 2). Individual low-dimensional inorganic nanostructures were mounted on thin metallic (Au, Al) wires (Ø 250 μm), which may be piezo-driven in three dimensions (with a precision better than 1 nm). The successful experimental runs include multi-walled BN (Figures. 1,2) [1-4] and N-doped C nanotubes (NTs) [5,6], ZnO nanorods and nanowires [7], double-walled CNTs, multi-walled CNTs filled with Cu [8], Au, CuI or ZnS, inorganic nanothermometers, e.g. liquid Ga-filled MgO NTs and In-filled SiO2 NTs, CdS nanobelts, BN cones and so on. The current-voltage (STM-TEM, Figure 1) and/or force-displacement (AFM-TEM, Figure 2) curves were recorded under a full control of the nanomaterial morphological, crystallographic and chemical changes/transformations during all stages of probing/manipulation. This allowed us to unambiguously interpret many new intriguing physical, chemical and electromechanical phenomena peculiar to novel advanced nanostructures. These include, but are not limited to deformation-driven electrical transport and piezoelectricity in individual BNNTs, time-resolved electrical transport in N-doped CNTs, rheostat-like behaviour, femtogram (Cu) and attogram (CuI) mass transports, and nanopipetting in Cu- and Cu-halide-filled CNTs. 1. 2. 3.
D. Golberg et al., Adv. Mater. 19 (2007), p. 2413. D. Golberg et al., Nano Lett. 7 (2007), p. 2146. D. Golberg et al., Acta Mater. 55 (2007), p. 1293.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 115–116, DOI: 10.1007/978-3-540-85226-1_58, © Springer-Verlag Berlin Heidelberg 2008
116 X.D. Bai, D. Golberg et al., Nano Lett. 7 (2007), p. 632. P.M.F.J. Costa, D. Golberg et al., Appl. Phys. A 90 (2008), p. 225. P.M.F.J. Costa, D. Golberg et al., Appl. Phys. Lett. 91 (2007), p. 2231081. P.M.F.J. Costa, D. Golberg et al., J. Mater. Sci. 43 (2008), p. 1460. D. Golberg et al., Adv. Mater. 19 (2007), p. 1937. The authors are indebted to Dr. O. Lourie and Mr. K. Kurashima for a continuous technical support.
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Figure 2. Force-displacement curves directly recorded on a relatively thick (a) and a thin (b) individual multi-walled BN nanotube using an AFM-TEM setup. The insets display the nanotube appearance at different stages of the experiments.
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Atomic structure of SW-CNTs: correlation with their growth mechanism and other electron diffraction studies R. Arenal 1, M.F. Fiawoo1, R. Fleurier1, M. Picher2, V. Jourdain2, A.M. Bonnot3, A. Loiseau1 1. LEM, CNRS-ONERA, 92322 Châtillon - France 2. LCVN, CNRS - Université Montpellier II, 34095 Montpellier - France 3. Institut Louis Néel CNRS/UJF, 38042 Grenoble - France [email protected] Keywords: Electron Diffraction, Atomic Structure, Growth Mechanism, Nanotubes, Nanofilaments
In this work, we have performed a comparative study of the helicity distribution in single wall carbon nanotubes (SW-CNTs) synthesized by different CVD processes in order to understand how the synthesis conditions can influence and control their growth mechanism [1]. To achieve this aim, we used the nano-beam (employing a parallel probe of about 50 nm) electron diffraction mode, which has been shown to be the most reliable technique for obtaining the complete determination of the atomic structure (chiral angle (θ) and diameter) of individual nanotubes (see for instance [2]). We recorded the intensities of electron diffraction patterns (EDPs) from individual SW-CNTs as well as from bundles of CNTs, using two different FEI Tecnai F20 TEMFEGs operating at 100 and 120 keV respectively. In the case of individual SW-NTs, their chiral indices were determined as the best value fitting experimental and simulated patterns calculated using kinematical electron theory. An example is shown in Figure 1 (a), which presents the EDPs of a SW-CNT, with in the inset the bright field image recorded under nanoprobe illumination of this tube. The structural parameters measured of this NT are θ=24.0 and diameter (extracted from the oscillations of the Eq-L) 2.8 nm. With these values, we simulate the corresponding EDPs and determined the chiral indices as the best fit between experimental and simulated EDPs. The SW-CNT of Figure 1 is thus unambiguously identified as a (24,17). In this contribution, we will present the comparative analysis of a significant number of NTs (>50) synthesized by two different CVD techniques: a thermal CVD process for which catalyst particles are prepared ex situ using a colloidal technique and deposited on the substrate and a hot filament CVD process where catalyst particles are prepared in situ by heating, under hydrogen gas, a few monolayers of the catalytic metal deposited on the substrate. In both cases, the substrates were perforated silica membranes, making possible to acquisition of EDPs from suspended nanotubes [1]. Furthermore, in this contribution we will also present our recent results on a very particular class of nano-filaments/NTs synthesized via CVD process. Figure 2 (a) shows a series of micrographs of a rope of 2 SW-CNTs and of their modification into a nanofilament. The analysis of the EDPs recorded on the rope section as well as on the filament section (Figure 2 (b) and (c), respectively) shows that we are dealing with two
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different kinds of objects: NTs (on the left of the TEM images) and their transformation to another structure (on the right) which keeps constant the NTs’ chiral angle but loss their transitional periodicity and their cylinder shape. Our ongoing experiments and analysis on these very interesting objects will be also presented. 1. 2.
M.-F. Fiawoo A.-M. Bonnot, V. Jourdain, R. Arenal, J. Thibault, A. Loiseau, J. Nanoscience and Nanotec, to be submitted (2008). R. Arenal, M. Kociak, D.J. Miller, A. Loiseau, Appl. Phys. Lett, 89, 073104 (06).
Figure 1. (a) and (b) experimental and simulated EDPs of a SW-CNT (inset: bright field image), respectively. We assign these EDPs unambiguously to a (24,17) NT.
Figure 2. (a) TEM micrographs showing a rope of 2 NTs and their transformation to a nano-filament. (b) and (c) EDPs recorded on the rope and on the nano-filament, respectively.
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TEM investigation ofSe nanostructures in/on Acetobacter xylinum cellulose gel-film N. Arkharova1, V.V. Klechkovskaya1, E. Suvorova1,2 1. A.V. Shubnikov Institute of Crystallography of Russian Academy of Sciences, Leninskii pr.,59, 119333, Moscow, Russia 2. Centre Interdisciplinaire de Microscopie Electronique, Ecole Polytechnique Federale de Lausanne, Station 12, CH-1015 Lausanne, Switzeland [email protected] Keywords: bacterial cellulose gel-film, polymer stabilised nanoparticles, selenium nanowires
Acetobacter xylinum cellulose (A X –BC) and its composites with some mineral and organic additives are the perspective materials for healing the heavy and complicated wounds. The Se/Ag/A X –BC composite is expected to have some antiphlogistic and immunostimulating properties [1]. This work reports the results of TEM/EDS analysis of the Se particles obtained from aqueous solutions by redox process in the presence of the N-poly (vinyl-2-pirrolidone) (PVP) and then absorbed by the surface and embedded into the AX BC gel-films. The PVP/Se ratio between 0 and 1 weight % was used. The solutions contained the traces of AgNO3. TEM images and SAED patterns, EDS data were processed and analyzed using Gatan Digital Micrograph 3.11.1, INCA (Oxford) and JEMS [2] software. TEM samples were the thin slices of the gel-film sections in epoxy. Two types of Se nanostructures were found: ellipsoidal particles in the range of sizes from 20 nm to 100 nm (Figure 1) and nanowires of a few nm width and up to a few microns long (Figure 2). Ellipsoidal particle were revealed both on the surface and ∼2 microns deep in the gel-film while the nanowires were only on the gel-film surface. X-ray EDS microanalysis and SAED patterns showed that ellipsoidal particles were the composites Se/ orthorhombic Ag2Se with different ratio between Se and Ag2Se. Nanowires had the structure of trigonal Se and appeared only in the solution with the Se concentration higher than 0.5% PVP/Se. 1. 2. 3.
Yu. G. Baklagina, A.K. Khripunov et al. Russ. J. Appl. Chem. V. 78 (2005) P. 1176-1181 P. Stadelmann, JEMS. http://cimewww.epfl.ch/. 2008. We kindly appreciate to our colleagues from Institute of Macromolecular Compounds, RAS, for given samples.
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Figure 1. TEM image of the Se/Ag/A X-BC composite cross-section in epoxy
Figure 2. TEM image of selenium nanowires
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In-situ electron irradiation studies of metal-carbon nanostructures L. Sun1, Y. Gan1, J.A. Rodriguez-Manzo2, M. Terrones3, A.V. Krasheninnikov4 and F. Banhart2 1. Institut für Physikalische Chemie, Universität Mainz, D-55099 Mainz, Germany 2. Institut de Physique et Chimie des Matériaux, UMR 7504, F-67034 Strasbourg 3. Advanced Materials Department, IPICyT, 78216 San Luis Potosí, México 4. Accelerator Laboratory, University of Helsinki, P.O.B. 43, FIN-00014, Finland [email protected] Keywords: nanotubes, carbon, in-situ electron microscopy
The properties and the behaviour of nanoparticles are subjects of highest current importance. Experiments on individual clusters are generally difficult but can be carried out by the techniques of modern in-situ electron microscopy. The electron beam can be used as a tool to induce structural changes on an almost atomic scale [1]. Particularly well-suited systems for electron irradiation studies are metal nanocrystals in graphitic shells. Under electron irradiation, the shells contract and lead to compressive forces. By measuring the lattice spacings in HRTEM images of metal crystals inside graphitic shells, pressures on the order 10-20 GPa were determined [2]. Nonhydrostatic pressure, e.g., inside collapsing nanotubes, may deform the crystals considerably. This can be made use of for studying the deformation behaviour of individual nanometer-sized crystals. Figure 1 shows how a Co wire inside a carbon nanotube is deformed and extruded when the nanotube collapses locally under electron irradiation [3]. Another deformation cell, allowing the detailed study of crystal deformation, was designed by electron-beam structuring, as shown in figure 2. A 'carbon onion', encapsulating a Au crystal, was punctured by a fully focused electron beam and subsequently exposed to uniform irradiation so that the Au crystal was slowly extruded through the hole. The experiment was carried out at different temperatures and with different metals, and the deformation of metal crystals was studied in detail [4]. Electron irradiation of metal-carbon nanocomposites leads to the ballistic displacement of atoms from the graphitic shells. Some carbon atoms are sputtered into the metal crystal and may, depending on the solubility and diffusivity of carbon in the respective metal, segregate at the metal surface. This is shown in figure 3 where carbon atoms are ingested into a FeCo crystal in a multi-wall carbon nanotube [5]. A new nanotube nucleates and grows through the channel of the host nanotube. This experiment allows us to study the nucleation and growth of nanotubes by in-situ electron microscopy. The interaction between metal and carbon was also studied on an atomic scale by introducing Au and Pt atoms into graphene layers. Individual Au or Pt atoms were observed by HRTEM at different temperatures, and their migration was monitored [6]. [7] 1.
A. Krasheninnikov and F. Banhart, Nature Materials 6 (2007) 723.
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2. 3. 4. 5. 6. 7.
L. Sun, J.A. Rodriguez-Manzo and F. Banhart, Appl. Phys. Lett. 89 (2006) 263104. L. Sun, F. Banhart, A. Krasheninnikov, J.A. Rodriguez-Manzo, M. Terrones and P.M. Ajayan, Science 312, (2006) 1199. L. Sun, A. Krasheninnikov, T. Ahlgren, K. Nordlund and F. Banhart, submitted. J.A. Rodriguez-Manzo, M. Terrones, H. Terrones, H.W. Kroto, L. Sun and F. Banhart, Nature Nanotech. 2 (2007) 307. Y. Gan, L. Sun and F. Banhart, Small, in the press. Support by the DFG (Ba 1884/4-1) is gratefully acknowledged.
Figure 1. Electron irradiation of a carbon nanotube encapsulating a Co crystal. The collapse of the tube leads to the extrusion of the metal [3].
Figure 2. Extrusion of a Au crystal through a focused-beam-induced hole (top left) from a spherical carbon shell under subsequent uniform electron irradiation [4].
Figure 3. Growth of a multi-wall carbon nanotube (arrowed) from a FeCo crystal inside a host nanotube under uniform electron irradiation [5].
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Application of 80kV Cs-corrected TEM for nanocarbon materials A. Chuvilin1, U. Kaiser1, D. Obergfell2, A. Khlobystov3, S. Roth2 1. University of Ulm, Central Facility of Electron Microscopy, Electron Microscopy Group of Materials Science, Albert Einstein Allee 11, 89069 Ulm, Germany. 2. Max Planck Institute for Solid State Research, Heisenbergstraße 1, 70569 Stuttgart, Germany. 3. School of Chemistry, University of Nottingham, University Park, Nottingham NG7 2RD, UK. [email protected] Keywords: TEM, Cs correction, carbon materials
Being one of the main methods for studying the atomic structure of carbon materials, high resolution TEM suffers from a serious flaw as the energy of electrons, necessary to obtain atomic resolution (200-300 keV), is far above the kick-off damage threshold for graphitic carbon (about 60 keV) [1], which means that carbon nanostructures are always observed with a structure more or less distorted by electron beam. Practical implementation of Cs correctors in the last decade [2] opened the possibility to decrease acceleration voltage of the microscopes significantly, yet preserving the resolution at reasonable level. Decreasing the energy of the electron beam, besides decreasing radiation damage, has the advantage of increasing the phase contrast [3]. Considering the possibility to visualize one single carbon atom, the increase of the contrast reduces quadratically the electron dose necessary for this. Here we report our first experience in application of 80kV electron microscope with imaging Cs corrector (Titan 80-300, FEI, Holland) for studying carbon nanotubes, peapods and graphene. We show that nano-structures composed of graphitic carbon are stable under the conditions of observation for more than 30 minutes, that opens the possibility for monitoring of single atom dynamics, diffraction studies and EELS measurements without radiation damage induced by the electron beam. Visualization of one single carbon atom within a graphene network is electron-optically achievable, but is limited in a single shot by the dynamical range of the CCD camera. Thus a multiple frame acquisition-averaging method is proved to be adequate for accumulation the electron dose necessary for resolving single carbon atom. Direct visualization of defects within the graphene network is demonstrated (Figure 1); in particular the stability of a mono-vacancy is shown under the beam irradiation. Though almost eliminating kick-off damage unavoidable at medium range electron energies, decreasing the energy to 80keV significantly increase ionization and heating of the sample. Ionization is not an issue for graphitic carbon materials due to high electrical conductivity, while special precautions should be taken against local overheat. From the other hand side the elevated temperature inside the beam spot opens possibilities for dynamical observation of heat activated chemical reactions and
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structure transformations. On the example of a heat induced transformation of SWNT peapod into double wall nanotube it is shown that good heat sink is essential for observation of pristine atomic structure of sensible materials at low voltages. At the conditions, when an area under observation is thermally isolated, heat induced evolution of the structure of SWNT metallofullerene peapods is observed, namely: breakage of the walls and coalescence of fullerene balls, behavior and interaction of individual metal atoms, inner tube formation from fullerenes, coalescence and collective action of Dy atoms resulting in SWNT breakage, self curing of the tube. At similar conditions the dynamics of Ag atoms on the surface of graphene is monitored. A formation of unusual 2D surface crystals with 4-fold symmetry is shown. Utilization of accelerating voltages below 100kV at Cs corrected TEM opens possibilities for a broad field of research of pristine atomic structure of materials including light elements like carbon and point defects like single vacancies. 1. 2. 3.
R.F. Egerton, P. Li, M. Malac, Micron 35 (2004), p. 399. M. Haider, H. Rose, S. Uhlemann, E. Schwan, B. Kabius, K. Urban, Ultramicroscopy 75 (1998), p. 53. M.T. Hayashida, T. Kawasaki, Y. Kimura, Y. Takai, Nuclear Inst. and Methods in Physics Research 248 (2006), p. 273
Figure 1 Monovacancy in a single graphene sheet. (a) Overview of graphene edge. The image is a drift corrected average of 10 frames. (b) Simulated image and (c) a model of monovacancy. Image is calculated for the conditions of the best visibility of the vacancy (Cs=-10μm, df=-8nm). (d) Simulated image considering electron dose of 106e-/nm2 (top) and experimental image obtained by single 1 sec exposition (bottom). (e) The same set as (d), but for the dose 107e-/nm2 and 10 frames average correspondingly.
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Control of gold surface diffusion on Si nanowires Martien .I. den Hertog1, Jean-Luc. Rouviere1, Florian Dhalluin2, 3, Pierre J. Desré3, Pascal Gentile4, Pierre Ferret3, Fabrice Oehler4, Thiery Baron2 1. Laboratoire d'Étude des Matériaux par Microscopie Avancée, CEA/DRFMC, CEAGrenoble, 17 rue des Martyrs, 38052 Grenoble Cedex 9 2. Laboratoire des Technologies de la Microélectronique, CNRS UMR 5129, CEAGrenoble 17 rue des Martyrs, 38052 Grenoble Cedex 9 3. Laboratoire de Photonique sur Silicium, CEA/LETI/DOPT, CEA-Grenoble 17 rue des Martyrs, 38052 Grenoble Cedex 9 4. Laboratoire Silicium Nanoélectronique Photonique et Structure, CEA/DRFMC, CEA-Grenoble, DRFMC, 17 rue des Martyrs, 38052 Grenoble Cedex 9 [email protected] Keywords: Silicon nanowires, STEM, Gold diffusion
Silicon nanowires (NW) were grown by the Vapour-Liquid-Solid (VLS) mechanism using gold as the catalyst and silane as the precursor. In this work we show that the diffusion of gold from the catalyst particle over the wire sidewalls can be controlled by adjusting growth parameters: partial silane pressure and growth temperature. The diameter of the NWs has also an effect. Gold diffusion results in the formation of gold clusters on the wire sidewalls (Figure 1). High Angle Annular Darkfield Scanning Transmission Electron Microscopy (HAADF STEM) and Scanning Electron Microscopy (SEM) were used to observe the presence or absence of gold clusters on the wire sidewall. In literature an influence of oxygen on the gold diffusion can be found [1]. Figure 2 illustrates the effect of growth temperature and partial silane pressure on the gold surface diffusion. Gold clusters are absent on the NW side walls for high silane partial pressure, low temperature and small NW diameters. The NW sidewall morphology is modified by the presence of gold. Different models will be qualitatively discussed in the presentation. The adsorption of silane on the NW sidewalls seems to be the main physical effect governing gold diffusion. 1. 2.
S. Kodambaka, J. B. Hannon, R. M. Tromp, F. M. Ross, Nano Lett. 6 (2006), 1292 The results reported in this publication were obtained with research funding from the European Community under the FP6 - Marie Curie Host Fellowships for Early Stage Research Training (EST) “CHEMTRONICS” Contract Number MEST-CT-2005-020513
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Figure 1. High resolution HAADF STEM image showing the gold clusters on the NW sidewall, the brighter particles are gold.
Figure 2. (A) SEM image of a wire grown at 500°C at low partial silane pressure (0.19 mbar) clearly covered by gold clusters and slightly tapered. (B) HAADF STEM image of a wire grown at 430°C at low partial silane pressure (0.19 mbar); gold clusters can be observed near the catalyst particle. (C) HAADF STEM image of a wire grown at 500°C at high partial silane pressure (1.023 mbar). Gold clusters can be observed on the 300 nm region close to the catalyst particle.
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Nanowires of Semiconducting Metal-oxides and their Functional Properties M. Ferroni1, C. Baratto1, E. Comini1, G. Faglia1, L. Ortolani2,3, V. Morandi2, S. Todros1, A. Vomiero1, and G. Sberveglieri1 1. INFM-CNR SENSOR and Dept. of Chemistry and Physics, Brescia University, via Valotti 9, 25133 Brescia, Italy 2. CNR-IMM Section of Bologna, via Gobetti 101, 40129 Bologna, Italy 3. University of Bologna, Dept. of Physics, v.le B. Pichat, 6/2, 40127 Bologna, Italy. [email protected] Keywords: Indium oxide, tin oxide, biosensor.
The increasing scientific interest in crystalline nanowires and nanorods has stimulated their functional exploitation, and single-crystalline nanowires are nowadays emerging as building blocks for a new generation of electronic [1,2,3], and optoelectronic nanometer-scaled devices [4,5]. The structural and functional properties of nanostructures of ionic metal oxides are attracting almost the same consideration addressed to nanowires of silicon or group III-V compounds. The synthesis of tin- and indium- oxides nanowires is based on thermal decomposition of precursor powders followed by vapor-solid (VS) or vapor-liquid-solid (VLS) growth. Such a process is very promising for nanostructure fabrication, due to its simplicity with respect to the top-down technology of silicon processing. Microstructural investigation of the nanowires is a key-feature for a thorough knowledge of the growth mechanisms, which can be assisted by either seeding particles or noble metal catalysts (see Figure 1), enhancing the capability to control the dimension and the shape of the nanostructures [6]. The chemical homogeneity and the crystalline arrangement were investigated by SEM and TEM. HRTEM and electron diffraction revealed the single-crystal arrangement and the atomically-sharp crystal facets for the termination of the nanowires. EDX and STEM were used for the investigation of impurities and local variations in the composition. Further improvement of the preparation methodology addresses the capability to fabricate heterogeneous structures and chemically non-homogeneous nanowires. As shown in Figure 2, the transport and condensation process was tailored to obtain a longitudinally-assembled heterostructure of In2O3 and SnO2 [7]. The nucleation and growth of single-crystalline SnO2 nanowire at the termination of an In2O3 nanowire was obtained through the prominent catalytic activity of gold nanoparticles. For the investigation of the electrical properties of nanowires, piezo-actuated manipulators were operated into the SEM to fabricate single-wire-based devices. Presently, nanowires of tin oxide are presently under exploitation for applications in the field of chemical sensing and DNA detection.
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1. 2. 3. 4. 5. 6. 7.
Y. Cui and C.M. Lieber, Science 291 (2006) p. 851. X.Duan et al, Nature 409 (2001) p. 66. Y. Huang et al, Nano Letters 2 (2002) p. 101. C.M. Lieber, MRS Bullettin 28 (2003) p. 486. L. Samuelson et al, Materials Today 6 (2003) p. 22. A. Vomiero Crystal Growth and Design 7 (2007) p. 2500. A. Vomiero et al., Nano Letters 7 (2007) p. 3553.
Figure 1. (left) Nucleation and growth of In2O3 nanowires from the In2O3 buffer layer over a sapphire substrate. (right) TEM image of the catalytic gold particle at the tapered apex of a crystalline In2O3 nanowire prepared by VLS.
Figure 2. Structural characterization of the longitudinal In2O3-SnO2 heterostructure. (main picture) TEM panoramic view of the In2O3 nanowire, with the SnO2 nanowire extending lengthwise. The black arrow marks the termination of the SnO2 nanowire and the catalytic Au nanoparticle, which assisted the VLS growth of heterostructure. (1) SAED pattern showing the cubic single-crystal arrangement for the indium oxide nanowire. (2) CBED pattern and High-Resolution image from the tin oxide nanowire, demonstrating its single-crystalline cassiterite arrangement. (3-4) highly magnified TEM view and corresponding SAED pattern of the heterojunction, where superimposition of both indium and tin oxides has been recorded.
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Phase relations in the Fe–Bi–O system under hydrothermal conditions A. Gajović1, 2, S. Šturm1, B. Jančar1, M. Čeh1 1. Jožef Stefan Institute, Jamova 39, SI-1000 Ljubljana, Slovenia, 2. Ruđer Bošković Institute, Bijenička 54, HR-1002 Zagreb, Croatia, [email protected] or [email protected] Keywords: hydrothermal synthesis, bismuth ferrite, hematite, goethite, bismuth oxide
The phases appearing during a hydrothermal reaction in the Fe–Bi–O system were investigated, with the aim to optimize the conditions for the syntheses of nanostructured bismuth ferrite. BiFeO3 is a multiferroic material; hence it is ferroelectric and antiferromagnetic at room temperature [1]. Since the spins in this material take the form of a long-wavelength (62 nm) the spiral linear magnetoelectric effect averages to zero. One of the ways to recover this linear effect is with thin-film epitaxial constraints. Thus, we expect the same effects in nanostructured BiFeO3. For the hydrothermal reactions two different precipitation procedures were carried out with various molar ratios (x) of iron and bismuth ions. In the first procedure the solutions containing both iron and bismuth ions (starting compounds Fe(NO3)3·9H2O and Bi(NO3)3), were co-precipitated with tetraethyl ammonium hydroxide (TMAH), while in the second procedure the precipitation of the iron and bismuth ions was carried out separately, prior to mixing in the autoclave. Before the precipitation the Bi(NO3)3 was dissolved in HNO3 by vigorous stirring. The hydrothermal treatments were performed in the autoclave at pH 13.54 at 200°C for 6 h. The stability of the phases was checked in the case of 10 mol% Bi(NO3)3 by one week of hydrothermal treatment. The prepared phases were analyzed by X-ray powder diffraction (XRD), while the morphologies and nanostructure of the different phases were determined using highresolution transmission electron microscopy (HRTEM) and/or scanning electron microscopy (SEM). The chemical compositions were examined by energy-dispersive Xray analysis (EDXS). The phases observed in the prepared samples are listed in Table I. The tailoring of the goethite morphology in both series of precipitations was investigated by TEM. It was observed that the aspect ratio of goethite nanorods increased with the increase of Bi3+ ions in the reaction. However, other phases, like hematite and bismuth ferrite, were also produced. We obtained pure bismuth ferrite only in a co-precipitated sample with x=50%. In the case of separate precipitation, BiFeO3 was produced in a reaction with a lower content of Bi3+. Again, various phases coexisted in the sample. Goethite and hematite were shown to be stable phases (Figure 2). In hematite, containing approximately 7.2% of Bi3+, stacking faults were observed (Figure 2b). These are potential sites for Bi3+ accommodation, thus indicating a possible mechanism for the synthesis of nanosized BiFeO3.
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W. Eerenstein, N. D. Mathur and J. F. Scott, Nature 442 (2006), p. 759. The authors acknowledge financial support from the European Union under the Framework 6 program under a contract for an Integrated Infrastructure Initiative. Reference 026019 ESTEEM.
Table I. The phases obtained with hydrothermal reactions based on the XRD analysis. x indicates mol% of Bi3+ in the reaction. x (mol %) Bi+3 2
5
10
20
30
50
Co-precip.
G*
G
G, H*
G
G, Bi25FeO40
G, BiFeO3, Bi2O3
G, H, BiFeO3 G, H BiFeO3, Bi2O3
BiFeO3
Separate precip.
amorph. phase G, H, BiFeO3, Bi2O3
H, BiFeO3, Bi2O3
10 (one week)
G, H
* G→goethite (α-FeO(OH)); H→hematite (α-Fe2O3)
a)
b)
Figure 1. a) Bismuth ferrite hydrothermally synthesized in a co-precipitated sample with x = 50 %. b) High-resolution micrograph of synthesized BiFeO3. a)
b) H
G
Figure 2. a) Goethite nanorods (G) and hematite (H) as stable phases after one week of the hydrothermal reaction. b) HRTEM of hematite with stacking faults (denoted by arrow).
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Dose dependent crystallographic structure of InAs nanowires F. Gramm1, E. Müller1, I. Shorubalko2, R. Leturcq2, A. Pfund2, R. Wepf 1 and K. Ensslin2 1. Electron Microscopy ETH Zurich – EMEZ, 2. Laboratory of Solid State Physics Swiss Federal Institute of Technology - ETHZ, CH-8093 Zurich, Switzerland [email protected] Keywords: InAs nanowires, transmission electron microscopy, crystal structure
Semiconductor technology has improved the procedures to miniaturize devices step by step. Usually these devices are produced by lithographic methods on planar substrates. Another approach is to grow nano-objects, which could then be assembled to an electronic device ‘bottom-up’ as discussed in [1,2]. The InAs-nanowires investigated in this study were grown by Metal-Organic Vapor Phase Epitaxy (MOVPE) on a GaAs 〈111〉B substrate. The wires were grown catalytically from gold nanoparticles, which were randomly distributed on the substrate, as it was first performed and described by Wagner and Ellis [3]. The crystallographic structure of semiconductor nanowires is one of the important properties, which determine the electronic or optical properties of these miniaturized devices. To explore the structural organisation, nanowires were investigated by HRTEM (FEI Tecnai F30 FEG; 300kV, point to point resolution 0.19nm). The InAs nanowires crystallise in zinc-blende or in wurzite structure, which are the two maximally closedpackings for InAs. Selected area electron diffraction (SAED) was used to determine which of the two structures was present in the analysed nanowires. The nanowires, grown as described by A. Pfund et al. [4], were found to have mainly wurzite structure as confirmed by high-resolution electron microscopy (HRTEM) (Figure 1). In additions a small approximately 3nm thick amorphous oxide layer covering the surface of the nanowires was found.. No defects in the crystal lattice were visible in the first exposure during HRTEM imaging. However on the following exposure (Figure 2), taken some minutes later from the same area of the nanowire, lattice fringes were also observed in the oxide layer. Further increase of the electron dose resulted in a more and more complete distortion of the nanowire. A patchwork of lattice fringes became visible in HRTEM images and at some positions of the nanowire even zinc-blende SAED-pattern as shown in Figure 2 were obtained, which could not be found under low dose exposure. It is well known that electrons interact strongly with matter and can therefore also change or damage samples. Here we show that the electron irradiation may change the structure of nanowires. From the SAED pattern it is also clear, that both wurzite and zinc-blende structure occur in these nanowires. Due to the fact that the SAED-pattern of the zinc-blende structure was only observed after extended irradiation, it may be that the
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zinc-blende type structure represents only transformed wurzite type structures from non irradiated nanowires. To be on safe side for structure investigations of e.g. sensitive nano-devices it will be necessary to record a pair of images - one at minimal dose and one at extended irradiation - to document beam induced material reorganisation. 1. 2. 3. 4.
C.M. Lieber, Sci. Am. 285 (2001), p.58. W. Lu, C.M. Lieber, Nature Materials 6 (2007), p. 841. R.S. Wagner, W.C. Ellis, Appl. Phys. Lett. 4 (1964), p. 89. A. Pfund et al., Chimia 60 (2006), p. 729.
Figure 1 HRTEM-image of InAs-nanowire. A) Nanowire cover with the small oxide layer. Some lattice fringes can be recognized in the oxide layer and a Moirée-pattern of two overlapping crystal lattices is visible in the nanowire. B) SAED-pattern of wurzite structure with viewing direction along 110. C) HRTEM-image of the InAs showing a patchwork of lattice fringes. D) The zinc-blende SAED-pattern which can be observed after some working time e.g. in such a distorted nanowire.
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HRTEM simulations of planar defects in ZnTe nanowires I. Häusler1, H. Kirmse1, W. Neumann1, S. Kret2, P. Dłużewski2, E. Janik2, G. Kraczewski2, and T. Wojtowicz2 1. Humboldt University of Berlin, Insitute of Physics, D-12489 Berlin, Germany 2. Institute of Physics PAS, Al. Lotników 32/46, 02-668 Warsaw, Poland [email protected] Keywords: nanowire, HRTEM simulation, vapour-liquid-solid growth
Semiconductor nanowires are freestanding one-dimensional objects with a length of some microns and a diameter of few tens of nanometres. Due to this unique shape the electronic states of the nanowires are quantized perpendicularly to the growth direction. Such semiconductor structures are favourable for numerous applications. However, semiconductor nanowires very often contain planar defects. In order to prevent defects during the growth process it is very important to have a thorough understanding of their nature and formation. In this work we have investigated the structure of ZnTe nanowires grown on (001) oriented GaAs substrate by vapour-liquidsolid (VLS) deposition. A liquid droplet of gold with a sphere size of about 70 nm acts as a catalyst and defines the diameter of the beneath growing nanostructure during the molecular beam epitaxy. The defects were analysed by high-resolution transmission electron microscopy (HRTEM). Figure 1 shows a bright-field TEM image of a separated ZnTe nanowire on a holey carbon film. For the defect analysis HRTEM images were interpreted by computer simulation. Figure 2 is a HRTEM image observed in a [01-1] zone axis orientation of the nanowire. The planar defects are visible by the characteristic lines perpendicular to the <111> growth direction. On the basis of the HRTEM images the planar defects can be identified as stacking faults with {111} stacking fault planes. In the field of view, the orientation of the dumbbells is flipping from one lamella to the next lamella. The mean thickness of the lamellae varies from two to six monolayers (MLs). It seems that one orientation is more favourable than the other. We assume that these lamellae were formed by rotational twinning, where a 180° rotation around the [111] nanowire axis takes place. For HRTEM image simulation we constructed a super cell containing a twin lamella with a thickness of three MLs (cf. Figure 3). Figure 4 shows the results of the HRTEM image simulation as a function of the specimen thickness and the defocus values. The comparison of the experimental HRTEM image (Figure 2) and the simulated image of the super cell containing the twin lamella (Figure 3) reveals a good agreement for specimen thicknesses between 90 Å and 110 Å and defocus values between -120 Å and -170 Å. Both images show the zigzag pattern of the flipping dumbbells. In addition, HRTEM simulation will be presented where the influence of the cylindrical shape of the nanowire as well as astigmatism and defocus variation were considered.
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Figure 1: Bright-field image of an isolated ZnTe nanowire on a holey carbon film.
Figure 2: HRTEM image of upper region of a ZnTe nanowire (cf. marked area in Fig. 1).
Figure 3: Super cell for HRTEM simulations
Figure4: Results of the HRTEM image simulation for the super cell shown in Fig. 3
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A universal method for determination of helicities present in unidirectional groupings of graphitic or graphitic-like tubular structures H. Jiang1, D.P. Brown2, A.G. Nasibulin1 and E.I. Kauppinen1,3 1. Laboratory of Physics and Center for New Materials, Helsinki University of Technology, P.O. Box 5100, 02150 Espoo, Finland 2. Canatu Ltd., Tekniikantie 21, 02150 Espoo, Finland 3. VTT Biotechnology, Technical Research Center of Finland, 02044 VTT, Finland [email protected] Keywords: electron diffraction, helicity, nanotubes
Since the recognition of carbon nanotubes by Iijima in 1991, nano-scale tubular structures, including graphitic carbon nanotubes (CNTs) and graphitic-like boron-nitride nanotubes (BNNTs), have generated tremendous research and engineering interest because of their unique helical structures as well as extraordinary properties. The physical properties of nanotubes sensitively depend on their helicities. Applicationoriented research on tubular structures thus requires reliable control over the growth process for selectively forming nanotubes with desired helicity. To better understand properties, develop practical applications and control synthesis, it is crucial to advance our capability for structural characterization of nanotubes. It has been widely recognized that electron diffraction is a powerful technique for structural characterization of nanotubes with little constrains in contrast to other tools like optical measurements. Electron diffraction analysis of individual single-walled CNTs [1, 2] or BNNTs [3] enables explicit determination of their atomic structures specified by so-called chiral indices (n,m). In reality, nanotubes often form bundles, or exist as coaxial multi-walled tubes, i.e. in general, in the form of unidirectional groupings, rather than as individual tubes. However, current solutions for structural characterization of such groupings are far from adequate and there is an urgent need for a reliable and quantitative method to characterize structural helicity properties in realistic samples. Here we introduce a universal method [4] for complete determination of helicities present in multi-walled or bundled carbon or boron-nitride nanotubes or their structural analogs. A critical dimension characteristic of the structural property of the atomic bond length in the structures is discerned from their electron diffraction patterns (EDPs), by which, all helicities present in the groupings of nanotubes are identified. The method is illustrated by using a simulated EDP of a single-walled CNT bundle (Figure 1) containing four tubes (13,2), (18,7), (20,6), (26,6). The diffraction pattern was calculated using the DIFFRACT program [5] by assuming a bundle tilt angle of 15° with respect to the electron beam. Eventually all helicities present in the bundle have been successfully resolved with an accuracy better than 0.2°.
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The method is further applied successfully for helicity analysis of two multi-walled BNNTs from their experimental electron diffraction patterns shown in Figure 2, which are reproduced with permission from A. Celik-Aktas et al, Applied Physics Letters, 86, 133110 (2005).Copyright 2005, American Institute of Physics. The results show precise details about the existing or missing helicities in the structures. A further application of the introduced method has been made to CNTs synthesized in a laser oven and has disclosed interesting structural phenomena. This will also be presented. 1. 2. 3. 4. 5. 6.
H. Jiang, A. G. Nasibulin, D. P. Brown and E. I. Kauppinen, Carbon, 45 (2007), p. 662. M. Gao, J.M. Zuo, R.D. Twesten, I. Petrov, L.A. Nagahara and R. Zhang, Appl. Phys. Lett. 82 (2003), p.2703 . R. Arenal, M. Kociak, A. Loiseau and D.-J. Miller, Appl. Phys. Lett. 89 (2006), p. 073104 H. Jiang, D. Brown, A.Nasibulin and E.I. Kauppinen, Phys. Rev. Lett. (2008), submitted. Ph. Lambin and A. Lucas, Phys. Rev. B 56 (1997), p.3571. We kindly acknowledge financial support from the Academy of Finland.
Figure 1. A simulated electron diffraction pattern of a SWCNT bundle containing four individual tubes (13, 2), (18, 7), (20, 6) and (26. 6).
Figure 2. Electron diffraction patterns from multiwalled BNNTs with (a) a narrow range of chiral angle distribution; and (b) a wide helicity distribution. The presence of the layer-line gap "A", "B" and "C" in (b) indicates deficiencies of certain helicities in the multiwalled nanotube, which has been evaluated.
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Microstructure of (112) GaAs nanorods grown by MBE E. Johnson1,2, S.A. Jensen1, L.P. Hansen1, C.B. Sørensen1 and J. Nygård1 1. Niels Bohr Institute, Nano Science Center, University of Copenhagen, Denmark 2. Department of Materials Research, Risø DTU, Roskilde, Denmark [email protected] Keywords: GaAs, Nanowires, TEM
Semiconductor nanorods of GaAs have been grown by MBE technique on (100) GaAs wafers. Au nanoparticles acting as nucleation and catalyst sites were made on the oxide free wafers by evaporation at 660 °C. After annealing for 10 minutes in an As backing flux, growth of the nanorods was carried out at 515 °C, and during final cooling of the samples to room temperature the As flux was maintained. Most of the nanorods were found to grow along the As-terminated (111)B directions typically with a heavily faulted zincblende structure or with larger perfect segments of wurtsite separating faulted zincblende segments [1]. This is the same growth and microstructure morphology observed for GaAs rods grown on (100) substrates using gold-particle-assisted metal-organic vapour-phase epitaxy [2]. However, in the present experiments a small fraction – about 5% – of the nanorods were found to grow in directions different from (111)B. Using a cleaved wafer it was possible from the knowledge of the wafer geometry to make a geometric analysis in the SEM of these rods which were found to grow along the (112)A directions (Figure 1). The (112) rods often appear to have special lancet-shaped blade-like microstructure where one part of the blade has a very smooth and straight edge. This part has wurtzite structure with the (11-20) direction parallel to the blade and rod axis. The other part has a more irregular and rugged edge and the crystallographic structure is zincblende with the (112) direction parallel to the blade and rod axis (Figure 2). In other cases the (112) rods would have a perfect wurtzite structure containing lamellae of zincblende a few monolayers in thickness parallel to the rod axis (Figure 2 and 3a). The zincblende parts of the rods in general contain more defects in form of intersecting stacking faults, twin lamellae and occasional appearance of wurtzite segments (Figure 2 and 3b). Minority formation of unusual nanorods may be due to local differences in the growth conditions either from defects in the wafer surface or as suggested by Wu et al. [3] as due to local strain variations. This should be studied further to elucidate details in the growth processes leading to occurrence of anisotropic crystal growth. 1. 2. 3. 4.
E. Johnson, M. Aagesen, C.B. Sørensen, J. Nygård, P.E. Lindelof, Proc. 16th Int. Micr. Cong. Sapporo, Japan, (2006), p. 1168. A.I. Persson, M.W. Larsson, S. Stenström, B.J. Ohlsson, L. Samuelson, L.R. Wallenberg, Nature Materials, 3 (2004), p. 667. Z.H. Wu, X. Mei, D. Kim, M. Blumin, H.E. Ruda, J.Q. Liu, K.L. Kavanagh, Appl. Phys. Lett. 83 (2003), p. 3368. This work is supported by the Danish Natural Science Research Council.
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Figure 1. GaAs nanorods grown at 515 °C by MBE on a cleaved [100] GaAs wafer. The drawing shows the projection of the (112) directions on to the [011] plane of observation. The (112) rod indicated by arrow shows a typical morphology where one side is straight and smooth and the other side is wavy and rough. Only about 5% of the rods are of (112) type but they are about 10 times longer than other rods in the growth and therefore dominate the SEM image visually.
Figure 2. (112) GaAs nanorod. The bottom part with the straight edge has wurtzite structure with a (11-20) direction along the wire axis and the top part with the rougher edge has faulted and twinned zincblende structure with a (112) direction along the wire axis.
a)
5 nm b)
5 nm Figure 3. GaAs nanorods with wire axis along (11-20) in the wurtzite structure (a) and wire axis along (112) in the zincblende structure (b). Both structures contain faults, twins and lamellae of the opposite structure. Wire axes are vertical.
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Structural characterization of ZnO nanorods grown on sapphire substrate by MOCVD P.-H. Jouneau1, M. Rosina2, G. Perillat1, 2, P. Ferret2 and G. Feuillet2 1. CEA - INAC, SP2M / LEMMA, 17 rue des Martyrs, 38048 Grenoble cedex 9, France 2. CEA - LETI, Minatec, 17 rue des Martyrs, 38048 Grenoble cedex 9, France [email protected] Keywords: ZnO, nanowire, nanorod
ZnO is a promising material for light emitting devices at room temperature up to the ultraviolet spectral region thanks to a band gap of 3.37 eV and a large exciton binding energy (≈ 60 meV). The growth of defect-free material and the mastering of doping are however two requirements to fully exploit this potential. In this perspective, the fabrication of ZnO nanorods by catalyst-free metallorganic chemical vapour deposition (MOCVD) has recently attracted a large interest: the growth of vertical, well-aligned nanorods with uniform lengths and diameters, and without structural defects, has been reported, even if their formation is not yet well understood [1, 2]. In this study, the microstructure and the morphology of ZnO nanorods obtained without catalyst by MOCVD on (0001)-oriented sapphire have been investigated by SEM, HR-TEM and HR-STEM in an attempt to better understand the growth mechanisms. They are obtained for substrate temperatures of 750°C-850°C, using diethylzinc and N2O as precursors. Scanning electron microscopy revealed a high density of nanorods (up to 109 cm-2), aligned perpendicularly to the surface with a unique crystallographic orientation (Figure 1). The growth starts with a 2D monocrystalline layer, but of uneven thickness and with some porosity, followed by the formation of islands terminated by large hexagonal pyramids. Nanorods appear mostly at the top of these pyramids, with a hexagonal section (corresponding to a growth axis along the c-direction) and most often a facetted tip. Length, diameter and density vary with the growth time, reactor total pressure and N2O partial pressure, but each given sample present a high shape and size uniformity (with typical wires diameter of about 150 nm and length of about 4 μm). TEM gives more insight on the crystalline structure. Filtered convergent beam electron diffraction patterns along a <0 1 -1 0> axes and HR-STEM images reveal a Zn polarity of the nanorods, i.e. a growth direction corresponding to the [0001] direction. No structural defects are visible in the nanorods, but the analysis of defects existing in the underlying layer (dislocations and planar defects such as antiphase boundaries) allows us to give some insights on the possible growth mechanisms. 1. 2.
U. Özgur, Ya.I. Alivov, C. Liu, A. Teke, M.A. Reshchikov, S. Dogan, V. Avrutin, S.-J. Cho and H.Morkoç, J. Appl. Phys. 98 (2005) 041301 I.C. Robin, B. Gauron, P. Ferret, C. Tavares, G. Feuillet, Le Si Dang, B. Gayral and J.M. Gérard, Appl. Phys. Lett. 91 (2007) 143120.
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Figure 1. SEM images of ZnO nanorods vertically aligned on a (0001) sapphire substrate, showing uniformity in diameters and lengths. Nanorods grow on top of hexagonal pyramids, with a uniform crystallographic orientation, and most often a faceted summit.
Figure 2. STEM images of a nanorod showing no structural defects. Filtered <0 1 -1 0> CBED patterns taken at various distances from the centre evidence the Zn polarity of the nanorod when compared with simulated patterns.
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Nucleation of Metal Clusters on Carbon Nanotubes X. Ke1, A. Felten2, D. Liang1, S. Bals1, J.J. Pireaux2, J. Ghijsen2, W. Drube3, M. Hecq4, C. Bittencourt4 and G. Van Tendeloo1 1. EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium 2. LISE, University of Namur, B-5000 Namur, Belgium 3. DESY, Hasylab, D-22603 Hamburg, Germany 4. LCIA, University of Mons-Hainaut, B-7000, Mons, Belgium [email protected] Keywords: Carbon Nanotubes, metal contact, chemical interaction
Electromigration in metal interconnects limits the number of devices in a single microprocessor. This limitation has already led to the replacement of Al as interconnect material by Cu, which has a higher current carrying capacity of the order of 106 A/cm2. Nevertheless, as the dimensions of the interconnect lines continue to shrink, electromigration and stress-induced migration phenomena remain reliability concerns for inlaid copper interconnects. CNT’s, capable of carrying currents of the order of 109 A/cm2, appear as a good candidate to replace state-of-the-art inlaid copper interconnects in microprocessors. A key step for integration of CNT’s in actual devices is the formation of stable and low-resistance ohmic contacts. Palladium appears as the most promising contact metal; elimination of the Schottky barrier and ballistic transmission of electrons has been already reported [1]. However, there is no obvious reason why Pd should give a smaller barrier than, say, Pt. Why is its electrical behaviour different? Does palladium react chemically with the CNT? Here we investigate the interaction between CNT’s and three different metals, namely platinum, palladium and gold, which are thermally evaporated onto CNT’s. High Resolution Transmission Electron Microscopy (HRTEM) was used to monitor the nucleation and evolution of the metal coatings on the CNT’s, whereas chemical interaction was monitored by X-ray photoemission spectroscopy (XPS). It was observed that during the onset of the evaporation, clusters are formed at the CNT surface (Figure 1). Electron tomography will be carried out to gain threedimensional information on how the CNT is coated with clusters. The cluster size increases by forming a quasi-continuous coating onto the CNT surface when the amount of metal evaporated on the surface is increased. Functionalization of the CNT surface leads to a better dispersion of the clusters and reduces the cluster size by creating more active sites for clusters to nucleate. Figure 2 shows a metal cluster on top of a CNT surface. The preserved structure of the graphene layers under the cluster suggests the absence of a mixed metal-carbon phase. Further information and confirmation can be obtained by XPS. If a chemical reaction is present at the interface, then the chemical environment of the atoms at the interface will show new features in the XPS spectra. The absence of new features near the metal core levels spectra (Figure 3) proves that no Pt-C, Pd-C or Au-C bonds are formed.
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The increase in the FWHM of the core-levels and the shift in the binding energy (Figure 3) will be discussed in terms of intrinsic initial electronic structure and final states effect.
Figure 1. TEM image of CNT’s coated with Pt. Au 4f
Intensity (arb. units)
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Figure 2. TEM image of CNT decorated with Pd clusters.
20 Å 10 Å 5Å 2Å 1Å 0.3 Å
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Figure 3. XPS spectra shows the evolution of the metal core levels by increasing amount of metal evaporated onto the CNT’s surface. 1. 2.
J. Tersoff Nature 23 (2003), p. 622 This work is supported by the Belgian Program on Interuniversity Attraction Pole (PAI 6/08) and by DESY and the European Commission under contract RII3-CT 2004-506008 (IASFS). JG is research associate of NFSR (Belgium).
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EDX and linescan modelling for core/shell GaN/AlGaN nanowire analysis L. Lari1, R.T. Murray1, T. Bullough1 P.R. Chalker1, C. Chèze2, L. Geelhaar3 and H. Riechert 2 1. Department of Engineering, University of Liverpool, Liverpool, L69 3GH, UK 2. Paul-Drude-Institut für Festkörperelektronik, 10117 Berlin, Germany 3. Namlab, 01187 Dresden, Germany [email protected] Keywords: EDX, nanowires, GaN, AlGaN
The growth of gallium nitride nanowires (NW) using nickel seeds to promote axial growth on sapphire substrates has previously been demonstrated [1]. Furthermore it has been shown that control of preferential axial or radial growth can be achieved by changing the group III:V flux during growth using either N-rich or Ga-rich conditions respectively. The next challenge is controlled growth of NW heterostructures incorporating both GaN and AlGaN. This achievement would open the way to development of nanoelectronic devices such as single wire III-nitride based field effect transistors. A critical issue is the control of Al incorporation in GaN/AlGaN radial NW heterostructures. This study reports the Al distribution analysis in such radial nanowires heterostructures grown by plasma assisted Molecular Beam Epitaxy on c-sapphire. The sample consists of a GaN NW core initially grown under N-rich conditions followed by an Al0.2Ga0.8N “shell” grown under Ga-rich conditions. NWs have diameters typically between 20 and 100 nm and lengths of around 1-1.5 μm. Compositional analyses were made using energy dispersive X-ray analysis (EDX) in a VG HB601UX FEG-STEM operating at 100 KV and equipped with a windowless Si(Li) EDX spectrometer. EDX spectra from point analysis were quantified using thin film approximation with calculated K factors using both the Zaluzec and Mott-Massey models [2]. Line scans profiles were obtained integrating selected energy windows centered in the Al and Ga K lines peaks with an appropriated choice of windows for background removal from the raw profiles. EDX point analyses on the NWs show an aluminium concentration, in the radial shell, which decreases as you go from the NW growth tip towards the NW root end at the substrate (Figure 1). Elemental net intensities from transversal EDX line scans across the NWs are interpreted using a geometric model of the EDX excitation volume within the core and shell region assuming the NW cross-section to be either hexagonal or radial (Figure 2). The sensitivity of the model to the AlGaN shell thickness, as it changes along the NW length, is discussed as well with the implications on the growth mechanism for these NW structures in group III-rich condition.
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L. Geelhaar, C. Chèze, W.M. Weber, R. Averbeck, H. Riechert, T. Kehagias, P. Komninou, G.P. Dimitrakopulos, and T. Karakostas, Appl. Phys. Lett. 91 (2007) 093113. D. C. Joy, A. D. Roming and J. I. Goldstein in “Principles of analytical electron microscopy” (Plenum Press, New York) (1986), p. 163. This work was supported by the EU contract MRTN-CT-2004-005583 (PARSEM) and the IST project NODE 015783.
(A)
(B)
Figure 1. (A) STEM bright field image of GaN/AlGaN core/Shell NW heterostructure. Results report Al at% from EDX point analysis quantification using Zaluzec (MottMassey) cross-section models. (B) Variation of Al at% across the NW diameter consistent with the core/shell structure
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Figure 2. (A) STEM bright field image of a typical NW tip the arrow reporting the position of the line scan corresponding to the experimental profile in (C and D). (B) NW core shell structure obtained from the best fitting with both radial and hexagonal Xsection models. (C, D) EDX and modelled profiles assuming different shell thickness.
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Mo6S9-xIx nanowires: structure studies by aberration corrected high resolution TEM and STEM V. Nicolosi1, J.N. Coleman2, D. Mihailovic3 and P. Nellist1 1. Department of Materials, University of Oxford, OX1 3PH, Oxford, UK 2. School of Physics, University of Dublin, Trinity College, D2, Dublin, Ireland 3. Josef Stefan Institute of Ljubljana, Jamova 39, 1000 Ljubljana, Slovenia [email protected] Keywords: nanowires, HRTEM, STEM
In the last few years carbon nanotubes have attracted a great deal of attention because of their interesting properties, including dependence of their chemical and electronic behaviour on their small dimension, high anisotropy and tube-like structure. Since nanotubes are difficult to process, one-dimensional inorganic nanowires have been receiving growing attention as a viable alternative to nanotubes. The most promising of these are the Mo6S9-xIx family. Easy fabrication [1], easy dispersability and processability [2], and uniformity in terms of metallic character [3] as well as diameter, makes Mo6S9-xIx nanowires one of the most promising one-dimensional materials. In order to fully understand their properties it is imperative to confirm the atomic structure of those mono-dimensional objects. Here we present a combined aberration corrected high resolution transmission electron microscopy (HRTEM) and scanning transmission electron microscopy (STEM) study of Mo6S9-xIx nanowires. Annular dark-field (ADF) STEM allows imaging of the individual one-dimensional molecules that aggregate to form bundles (Figure 1), but is unable to image the lower atomic number S atoms. Phase contrast HRTEM is less sensitive to Z, and reveals complementary information. Electron microscopy, crystallographic and electron microscopy-simulated studies allow us to determine both the unit cell structure and the packing structure of the wires in the bundles. Studies on Individual Mo6S3I6 nanowires can be best described as one-dimensional molybdenumchalcogenide-halide clusters. Mo-octahedra appear surrounded by iodine atoms and connected by bridging planes of three sulfur atoms (Figure 2). The nanowires tend to be weakly bounded into bundles by van der Waals forces, in a triclinic lattice arrangement, with a nanowire to nanowire distance of 0.958nm [3]. Data for other nanowire stoichiometries, such as Mo6S4.5I4.5 and Mo6S2I8, will be also presented. 1. 2. 3.
D. Vrbanic, M. Remskar, A. Jesih, A. Mrzel; P. Umek, M. Ponikvar, B. Jancar, A. Meden, B. Novosel, S. Pejovnik, P. Venturini, J. N. Coleman, D. Mihailovic, Nanotechnology 15 (2004), 635-638. V. Nicolosi, D. Vrbanic, A. Mrzel, J. McCauley, S. O’Flaherty, D. Mihailovic, W. J. Blau, J. N. Coleman, Journal of Physical Chemistry B, 109 (2005), 7124-7133 V. Nicolosi, P. D Nellist, S. Sanvito, E. C. Cosgriff, S. Krishnamurthy, W. J. Blau, M. L. H. Green, D. Vengust, D. Dvorsek, D. Mihailovic, G. Compagnini, J. Sloan, V. Stolojan, J. D Carey, S. J Pennycook, J. N. Coleman, Advanced Materials, 19, 4, (2006) 543-547
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Figure 1. Typical Annular Dark Field STEM image of a MoSI nanowire bundle. The coexistence of four different patterns can be explained by the fact that, seeing that those systems possess very low sheer modulus, a twisting process along the longitudinal axis of the same nanowire bundle is very likely to happen. This therefore produces different projections. Image taken using a using a VG HB501 instrument fitted with a Nion Cs corrector.
Figure 2. Structure of a Mo6S3I6 nanowire. Red, purple and yellow represent in order molybdenum, iodine and sulfur atoms. The brackets define the nanowire unit cell.
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Discrete Atom Imaging in Carbon Nanotubes and Peapods Using Cs-Corrected TEM Operated at 100keV Luca Ortolani1,2,3, Florent Houdellier3,Marc Monthioux3 1. CNR-IMM Bologna, v. Gobetti, 101, 40129 Bologna, Italy 2. Physics Dept., University of Bologna, v.le B. Pichat 6/2, 40127 Bologna, Italy. 3. CNRS CEMES, 29, rue Jeanne Marvig, 31055 Toulouse [email protected] Keywords: Cs Corrected TEM, Peapods, Carbon Nanotubes
The filling of carbon nanotubes (CNTs) with inorganic crystal as well as organic molecules or fullerenes has been widely investigated over the last years [1]. Filled Multi-Walled CNTs and, first of all, Single-Walled CNTs have become the most promising hybrid-nanomaterials for many applications because of the strong interaction between the outer carbon shell and the inner material and the imposed 1D dimension, capable to modify the atomic structure of both the filling and the nanotube [2]. Structural characterization by TEM is hampered by SWCNTs very small dimension and by its relative instability under the electron beam. Notwithstanding the outer carbon shell may act as a protecting environment, allowing the observation of the structure and of the dynamic behaviour of the contained molecules, in conventional 200 keV imaging, beam electrons are capable of displacing carbon atoms from the graphene lattice, resulting in a rapid destruction of the tube. Moreover, in conventional TEM observations of filled SWCNTs, the nanotubes themselves have been regarded as almost contrast-less container materials, due to the limited spatial resolution [3,4]. The use of aberration-corrected TEMs and lower observation energies seems to be the most promising experimental approach to study this kind of materials. We report here the results obtained on carbon structures and radiation-sensitive materials with the SACTEM-Toulouse, a TECNAI F20 equipped with a spherical aberration (Cs) corrector that was able to be operated at 100kV, for the first time we believe. The new lens set-up allows easy imaging of CNTs, as can be seen in Figure 1a. The Cs, and all the aberrations up to the fourth order, has been completely removed, and the graphene lattice of the CNT can be resolved. Moreover, the corrector is capable also to set the spherical aberration to a specified negative value in order to compensate the residual resolution-limiting aberration, which is, in our set-up, the 8 mm fifth-order spherical aberration. Therefore, a negative Cs of -19 μm can be used to compensate for this aberration, the same way defocus can be used to compensate Cs [5,6], resulting in a clearer resolution of fullerenes structure, as shown in Figure 1b. An interesting feature of Cs-corrected TEM is that a small change in focus affects dramatically the contrast of different z-planes in the specimen [3]. This effect is evident in the focus series of Figure 2: in a) and c) the contrast of the contained fullerenes is enhanced; in b), the contrast of the graphene lattice is maximized, while fullerenes have almost disappeared.
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1. 2. 3. 4. 5. 6.
M. Monthioux, Carbon 40 (2002), p. 1809. J. Sloan et al., Acc. Chem. Res. 35 (2002), p. 1054. K. Hirahara et al., Nano Lett. 6 (2006) p. 1778 Y. Sato et al., Nano Lett. 7 (2007), p. 3704. O. Scherzer, J. Appl. Phys. 20 (1949), p. 20. S. Uhlehmann, M. Haider, Ultramicroscopy 72 (1998), p. 109.
Figure 1. a) Image of a SWCNT, using Cs=0. Graphene lattice is clearly resolved. b) Image of a peapod, Cs=-19μm.
Figure 2. Focus series of a peapod (Cs=0). Focussing is decreasing from a) (overfocus) to c) (underfocus).
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Extended Defects in Semiconductor Nanowires Peter Pongratz1, Youn-Joo Hyun2, Alois Lugstein2, Aaron Andrews2 and Emmerich Bertagnolli2 1. Institute of Solid State Physics,TU-Wien, Wiedner Hauptstrasse 8-10, A-1040 Vienna, Austria 2. Institute of Solid State Electronics, TU Wien, Floragasse 7, A-1040 Vienna , Austria [email protected] Keywords: nanowires, TEM analysis, extended defects
Epitaxial growth of semiconductor nanostructures has opened new frontiers in semiconductor physics. Silicon and various other III-V semiconductor nanostructures grown on silicon wafers with various orientations and under different growth conditions have been successfully demonstrated by many groups [1,2,3]. TEM and electron diffraction techniques which are able to analyse these whiskers in detail have shown that even unconventional hexagonal structures of Si, GaAs and Ge could be generated under these conditions. We have studied these nanowires changing nucleation and growth conditions ( e.g. gas fluences or pressures, temperatures etc.) and found not only perfect single crystalline structures but also extended defects such as dislocations, stacking faults, twins, and grain boundary and interface structures. Using TEM analysis and refined cross-section preparation techniques they can be seen along the whiskers which are related to these growth conditions [4]. It is clear that it is essential to understand the relations between defects and growth conditions if a perfect control of the nanostructure growth process is desired. It is clear that not only the growth directions are of importance but also the structure of the side walls, the nucleation conditions at the substrate and the wetting conditions at the catalyst particles’ interface in the case of VLS growth . We will present our results related to the branching of nanowire trees on a trunk of silicon or the regularity of the cross-sections in the case of heteroepitaxial tree-structures. We have analysed kinks, dislocations and microtwins related to changing growth conditions but also complicated stacking faults and facets related to anisotropy effects and variable surface orientations . Elastic strains and surface energies are certainly more important for the properties of these nanowires than for their extended three dimensional counterparts. We will present various examples how HRTEM electron microscopy and analysis are extremely valuable to find out these details on a nanometer scale.
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Figure 1 (a) Cross sectional TEM and HRTEM image of a GaAs whisker and the respective SAED pattern recorded along the nanowhisker axis. (b) HRTEM image of the GaAs whisker with the GaAs whisker tip in the inset. 1. 2. 3. 4. 5.
T.Martenson ,C. Svensson , B Wacaser et al. Nano Lett. 4 (1987) Björk,M, B. Ohlsson T. Sass , A. Persson et al. Nano Lett. 2, 87 (2002) H. Jagannathan, M. Deal, Y. Nishi et.al. J.Appl. Phys.100, 024318 (2006) A. Lugstein A Andrews, M. Steinmair, Y. Hyun et al. Nanotechnology 18, 355306 (2007) We kindly acknowledge support by Austrian Science Fund P-18080-N07; Sixth EU Framework (NoE) SANDiE ; TU Vienna-USTEM and Austrian Society for Micro and Nanoelectronics
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Surface chemistry along a single silicon nanowire: Quantitative x-ray photoelectron emission microscopy (XPEEM) of the metal catalyst diffusion O. Renault, A. Bailly, P. Gentile, N. Pauc, T. Baron, L.–F. Zagonel, N. Barrett 1. CEA-LETI, MINATEC, 17, rue des Martyrs, 38054 Grenoble Cedex 09, France. 2. CEA-DSM/INAC/SP2M/SiNaPS, 17 rue des Martyrs, 38054 Grenoble cedex 9, France. 3. Laboratoire des Technologies de la Microélectronique, UMR CNRS 5129, 17 rue des Martyrs, 38054 Grenoble cedex 9, France. 4. CEA-DSM/IRAMIS/SPCSI, CEA Saclay, 91191 Gif-sur-Yvette, France. [email protected] Keywords: silicon, nanowire, diffusion, XPEEM, threshold photoemission.
Understanding and controlling the growth, usually performed by Vapour-LiquidSolid method, of silicon nanowires with grown metal catalysts is crucial for optimized morphologies, surface and electronic properties of these nanostructures regarding their future applications. One of the issues is the suppression of the tapering effect resulting in a cone-shaped wire, and arising from the diffusion, during growth, of the metal catalyst along the sidewall. Indirect evidences for such a phenomenon and for the presence of gold on the sidewall have been proposed previously [1], before a qualitative and direct identification using SEM-EDX and TEM was reported for technologically relevant Si NWs [2]. However, such techniques cannot provide quantitative and surfacesensitive information relevant to the diffusion process, and moreover can be extremely localized, preventing to obtain a global picture of the surface chemistry at the scale of the wire. For the first time, we quantify at a mesoscopic scale the surface chemistry of single, tapered silicon nanowire (100-250 nm diameter), using a novel X-ray photoelectron emission microscopy (XPEEM) with aberration-corrected energy-filtering and soft xray synchrotron illumination (ID08 beamline, ESRF) [3]. This spectromicroscope enables full-field, energy-filtered imaging using either secondary or core-level photoelectrons, giving information respectively about the local work function or the chemical state of the considered elements [4]. With XPEEM we provide a direct, nondestructive chemical confirmation of gold migration along the sidewall surface from threshold and core-level photoemission localized nanospectra. In particular, the secondary electron nanospectra, generated over a typical 200x200 nm2 area at the photoemission threshold from the images shown in Figure 1, have a remarkable doublestructure indicative of two distinct chemical phases present at the surface; the work function of which can be retrieved by modeling using Henke’s functions [3]. The values of work-functions Φ1 and Φ2, close to 4.1 and 5.0 eV, respectively, are in agreement with reported values for silicon nanowires and gold. The presence of gold all along the sidewall is further confirmed with the Au4f7/2 core-level image and corresponding
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nanospectra (Figure 2). From the Au4f7/2 and Si 2p core-level local intensities weighed by the corresponding photoionization cross-sections, we have access to the partial coverage, θAu, of gold on the nanowire sidewall. Gold has an evenly distributed island structure, covering 0.42 ± 0.06 of the nanowire surface corresponding to ~1.8 ML, in full agreement with independent transmission electron microscopy results. 1. 2. 3. 4. 5.
J. B. Hannon, S. Kodamka, F. M. Ross and R. M. Tromp, Nature 440, 69-71 (20006). F. Dhalluin et al., J. Appl. Phys. 102, 094906 (2007). A. Bailly, O. Renault et al., Phys. Rev. Lett., submitted. M. Escher et al., J. Phys. : Cond. Matter 17, S1329 (2005). This work was carried out with the financial support of the French ANR “XPEEM” project 05-NANO-065-02. Special thanks go to the staff of the ID08 beamline at the ESRF for their technical help during the experiments.
Figure 1. Energy-filtered XPEEM images at the photoemission threshold (workfunction contrast) and nanospectra on the sidewall (I, II), the catalyst (III) and the gold substrate (IV). Contrast inversion occurs when the electron energy exceeds the work function of the gold substrate (5.1 eV). Φ1 and Φ2 correspond to the first and second photoemission threshold, respectively at low and high energy.
Figure 2. XPEEM Au4f7/2 core-level image (left) and nanospectra (center). Right : Gold coverage derived from the Au4f and Si 2p local core-level intensities. Schematic drawing of the main corridors through which automobilists can reach the Eurogress Centre. Drawing is not to scale.
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TEM characterization of metallic Ni nanoshells grown on gold nanorods and on carbon nanotubes J.B. Rodríguez-González, M. Grzelczak, M.A. Correa-Duarte, 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, nickel, shells.
Anisotropic, magnetic nickel shells were grown on gold nanorods [1] and on carbon nanotubes (CNT) [2] using Pt nanocrystals as catalysts. We used high resolution transmission electron microscopy (HRTEM) to characterize these well-defined metallic nickel shells. In gold nanorods, Pt nanocrystals were selectively grown on the nanorod tips [3], and Ni2+ was subsequently reduced over them, in a quasi-epitaxial growth demonstrated by HRTEM and electron diffraction. Figure 1a shows a distinct contrast between the darker gold core and the lighter Ni shell, with a high degree of roughness in the shell, which might indicate the presence of an outer, thin NiO shell. The area corresponding to the gold core is almost completely covered with Moiré fringes arising from the presence of two superimposed different crystalline lattices in a quasi-epitaxial growth, confirmed by nanobeam diffraction (NBD). The distribution of Au and Ni within the particles was also investigated by scanning transmission electron microscopy STEM-XEDS elemental mapping (Figure 1b). A detailed analysis of the interface between gold core and nickel shell (Figure 1c and 1d) shows the presence of dislocations in this interface. TEM images of the CNT/Pt@Ni samples (Figure 2a) evidence the homogenous coating of individual CNTs and reveal the polycrystalline nature of the Ni layer on the final magnetic nanowires (Figure 2b). STEM-XEDS elemental mapping (Figure 2c) show the relative distribution of the elements; clearly Pt is located in the inner side of the metallic shell while Ni is mostly covering the outer part, as expected for an onionlike structure. HRTEM and SAED analysis also confirm the presence of unchanged CNTs in the inside part. 1. 2. 3.
M. Grzelczak, B. Rodríguez-González, J. Pérez-Juste, and L. M. Liz-Marzán, Adv. Mater. 19 (2007), 2262. M. Grzelczak, M. A. Correa-Duarte, V. Salgueiriño-Maceira, B. Rodríguez-González, J. Rivas, and L. M. Liz-Marzán, Angew. Chem. Int. Ed. 46 (2007), 7026. M. Grzelczak, J. Pérez-Juste, B. Rodríguez-González and L.M. Liz-Marzán, J. Mater. Chem. 16 (2006), 3946.
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Figure 1. Nickel coated gold nanorods. a) Note the presence of Moiré fringes and the shell roughness. b) STEM-XEDS elemental mapping. c) Interface between gold core and nickel shell. d) FFT filtering of (c).
Figure 2. Nickel coated carbon nanotube. a) Bright field image of one CNT/Pt@Ni. b) Outer Ni layer. c) STEM-XEDS elemental mapping.
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Electron Irradiation Effects in Carbon Nanostructures: Surface Reconstruction, Extreme Compression, Nanotube Growth and Morphology Manipulation M. Terrones1, L. Sun2, J.A. Rodriguez-Manzo3, H. Terrones1, and F. Banhart3 1. Advanced Materials Department, IPICyT, 78216 San Luis Potosí, México 2. Institut für Physikalische Chemie, Universität Mainz, D-55099 Mainz, Germany 3. Institut de Physique et Chimie des Matériaux, UMR 7504, F-67034 Strasbourg [email protected],mx Keywords: electron irradiation, nanotube growth, extreme compression
It will be demonstrated, that irradiation exposure at elevated temperatures, can be used as an effective tool to covalently weld SWNTs in order to create molecular junctions of various geometries [1-3]. We have fabricated Y, X and T-like junctions that are stable [2]. Tight binding molecular dynamics calculations demonstrate that vacancies, formed under the electron beam, trigger the formation of molecular junctions involving seven or eight membered carbon rings. We envisage that these results will pave the way towards controlled fabrication of novel nanotube-based molecular circuits, nanotube fabrics and network architectures. In this context, novel super architectures, using carbon nanotubes as building blocks will be discussed, and their mechanical and electronic properties presented, as well as their possible applications [4]. We will also show that the melting and solidification behavior of metal crystals can be drastically altered when they are encapsulated in fullerene-like graphitic shells [5]. The melting temperature of low melting point metal crystals (e.g. Bi, Sn, Pb, etc.) inside graphitic shells is increased relative to the bulk melting point by a much larger amount than that observed for metal crystals embedded in other materials. It appears that graphite is the ultimate material for enhancing the melting/solidification hysteresis of small crystals or clusters. Therefore, metal clusters encapsulated by graphitic shells may be potentially advantageous in temperature-resistant crystalline composite materials. In addition, we will demonstrate that controlled irradiation of MWNTs can cause large pressure buildup within the nanotube cores, to the extent of being able to plastically deform, extrude, and break solid materials that are encapsulated inside [6]. We further show by atomistic simulations that the internal pressure inside nanotubes can reach values higher than 40 GPa. Nanotubes can thus be used as robust nanoscale jigs for extruding hard nanomaterials and modifying their properties, as well as templates for other highpressure applications at the nanoscale. Finally, it was observed that when MWNTs containing metal particle cores are electron irradiated, carbon from graphitic shells surrounding the metal particles is ingested into the body of the particle and subsequently emerges as single- or multi-walled nanotubes inside the host nanotubes. These observations at atomic resolution in an electron microscope and indicate that bulk diffusion of carbon through the body of catalytic particles is the growth-limiting process [7].
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1. 2. 3. 4. 5. 6. 7.
M. Terrones, et al. Science 288 (2000), 1226. F. Banhart, et al. Small 1 (2005), 953. M. Terrones, et al. Phys. Rev. Lett. 89 (2002), 075505. J.M. Romo-Herrera, et al. Nano Letters 7, (2007), 570. F. Banhart, F., et al. Phys. Rev. Lett. 90 (2003), 185502. L. Sun, et al. Science 312 (2006), 1199. J.A. Rodríguez-Manzo, et al. Nature Nanotechnology 2 (2007), 307.
Figure 1. SWNT growth from a Co crystal. The growth happens under electron irradiation (intensity 100 A cm-2) at a specimen temperature of 600 °C. a) Image before growth; b) Image after 345 s of irradiation: and c) Image after 475 s.
Figure 2. (a-c) HRTEM images of a “T-like” junction formed after irradiating a “Y” junction. It is possible to thin one of the tubes of the “Y” junction, which eventually breaks and forms a “hook” (not shown here). It is also possible to observe the evolution of this junction during irradiation and the rotation by 180° under the electron beam; note the circular cross-section in one of the tubes (b).
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Crystallographic phase and orientation analysis of GaAs nanowires by ESEM, EDS, TEM, HRTEM and SAED A.M. Tonejc1, S. Gradečak2, A. Tonejc1 , M. Bijelić1, H. Posilović3, V. Bermanec3 and M. Tambe 2 1. Faculty of Science, Dept. of Physics, Bijenička 32, 10002 Zagreb, Croatia 2. MIT, Dept. Mat. Sci. Engng., 77 Massachusetts Avenue, MA02139 Cambridge, USA 3. Faculty of Science, Dept. of Geology, Horvatovac 102, 10002 Zagreb, Croatia [email protected] Keywords: GaAs nanowires, ESEM, TEM/SAED
The novel properties of semiconductor nanowires are interesting [1] for application and could be useful for application in nanoelectronics and photonics. Depending on the band gap and size of particular semiconductor nanowires, the shift in absorption/emission is observed. The samples were grown by general synthetic method developed by the C.M. Lieber group [2]. The method controls both the diameter and the length of nanowires (NW) during growth. This approach uses monodisperse nanocluster catalysts Au to define both the nanowire diameter and initiation of nanowire elongation during growth by a vapour–liquid-solid mechanism. In this article, as prepared GaAs NW samples, grown on GaAs [100] and [111] oriented films were synthesized using gold nanocluster catalyst and metalorganic chemical vapour deposition. Figure 1a shows ESEM photographs of NW vertically grown on GaAs [111] oriented film. It is the region from the edge of the sample to measure distribution of the NW lengths (Figure 1b) and the distribution of catalytic gold particles displayed in Figure 1c. The diameter of the catalytic particles at the top of the NW of diameter is (94±20) nm, that had to be the same as NW diameter. Belonging EDS spectrum (Figure 1d) shows the presence of Ga, As and Au. Figure 2a shows NW of GaAs observed by ESEM and EDS grown on [111] films. In Figure 2c TEM image of NW having diameter d= (32 ±10) nm, measured from TEM, is shown. The corresponding SAED in [111] Au orientation of the catalytic particles shows as well as HCP the {001} GaAs spots. The SAED (Figure 2d) gives evidence that the substrate film had the same orientation. The growth direction of NW is [011] and the structure is identified as HCP GaAs, according to ICDD-PDF No. 80-0003. The distribution of the lengths of nanowires is given from 500 to 3500 nm with maximum values 1300 and 2000 nm, for two observed samples, Figure 1c and Figure 2b, respectively. From HRTEM images it is found that NWs are coated with 3 to 5 nm thick amorphous layer. 1. 2.
M. Law, J. Goldberg and P. Yang, Annu. Rev. Mater. Res. 34 (2004), p. 83. Y. Cui, L.J. Lauhon, M.S. Gudiksen, J. Wang and C.M. Lieber, Applied Physics Letters 78 (2001), p. 2214.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 157–158, DOI: 10.1007/978-3-540-85226-1_79, © Springer-Verlag Berlin Heidelberg 2008
158 V a ri a b l e : V a r2 , D i s t ri b u t i o n : L o g -n o rm a l
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Figure 2. Sample 2: (a) ESEM; (b) The distribution of the lengths measured by ESEM; (c) TEM measurements; (d) SAED in [111] Au orientation gives direction of the NW growth in the image (c).
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3-Dimensional Morphology of GaP-GaAs nanowires M.A. Verheijen1, R. Algra2, M.T. Borgström1, G. Immink1, E. Sourty4, L.F. Feiner1, W.J.P. van Enckevort 3, E. Vlieg3 and E.P.A.M. Bakkers2 1. Philips Research, High Tech Campus 11, 5656AE Eindhoven, the Netherlands 2. Nederlands Institute for Material Research (NIMR), 2628CD Delft, The Netherlands 3. Radboud University Nijmegen, Toernooiveld 1, 6525ED Nijmegen, The Netherlands 4. FEI Company, Achtseweg Noord 5, AAE, 5600 KA Eindhoven, the Netherlands [email protected] Keywords: nanowires, tomography, morphology
Semiconductor nanowires are promising candidates for enabling integration of new functionalities. based on the advantageous properties of III-V semiconductors, such as high carrier mobility and optical activity into existing silicon technology. Because of the small dimensions and the consequently large ratio of surface to bulk atoms, the nanowire surface morphology and resulting chemistry can considerably affect the nanowire (opto-) electronic properties, such as carrier mobility and luminescence quantum yield. An elegant way to suppress such surface related effects is to cap the wires with a wide bandgap material to form core/shell nanowire structures. So far, only the vapour-liquid-solid (VLS), i.e. the axial growth mechanism and the zinc blende crystal structure has been considered in relation to the of III-V nanowire side faceting. In order to design and optimise nanowire properties for optical, electrical or mechanical performance, it is essential to take control over the nanowire surface morphology by fully understanding the parameters affecting radial growth. We have investigated the morphology of heterostructured GaP-GaAs nanowires grown by metal-organic vapor phase epitaxy as a function of growth temperature and V/III precursor ratio. The vertical growth of these wires is accompanied by lateral growth (i.e. thickening of the nanowires). The ratio vertical-to-lateral growth is strongly dependent on the growth parameters. The study of heterostructured core-shell nanowires with transmission electron microscopy (TEM) tomography (Figure 1) allowed to resolve the three-dimensional morphology, and to discriminate between the effect of axial (core) and radial (shell) growth on the morphology. Combining tomography with HRTEM (Figure 3), SEM and HAADF, a temperature and precursor dependent structure diagram (Figure 2a) for the GaP nanowire core morphology was constituted, and the evolution of the different types of side facets during GaAs and GaP shell growth was characterised. An additional parameter determining the sidewall morphology is the density of stacking faults (SF). In low SF-density <111> grown wires, the most stable sidewall morphology is allowed to form, contrary to high SF-density <111> grown wires. STEM/HAADF and SEM imaging was used to characterise this sidewall morphology (Figure 2b,c).
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Figure 1. 3-dimensional reconstruction of the morphology of the bottom part of a GaPGaAs heterostuctured nanowire (obtained using STEM/HAADF tomography). 2’’
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Figure 2a: Schematic 3D representation of the various VLS GaP core morphologies. Morphology types 1, 2’’, 2, 2’ represent zincblende crystal structures, while type 3 shows the morphology of alternating zincblende (ZB) and wurtzite (WZ) segments in high temperature grown wires. HAADF image of low SF-density nanowire (b) and a transition area from low to high SF-density (c).
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Figure 3: High resolution TEM image of the various facets of ZB and WZ segments in a high temperature grown GaP nanowire.
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Characteristics of Indium-Catalyzed Si Nanowires Z.W. Wang1, Z.Y. Li1 and F. Iacopi2 1. Nanoscale Physics Research Laboratory, School of Physics and Astronomy, University of Birmingham, Birmingham B15 2TT, UK. 2. IMEC, Kapeldreef 75, B-3001 Leuven, Belgium [email protected] Keywords: Si nanowire, indium, structure
Si nanowires are promising candidates for applications in nanoelectronics, opoelectronics, and sensors. Au is a highly efficient catalyst that is widely used in the growth of Si nanowires by the well-known vapor-liquid-solid growth mechanism. However, the mid-gap trap states in Si introduced by surface or bulk diffusion of Au can be a strong killer to minor carrier lifetime, thus limit the use of Au-catalyzed nanowires in Si-based technology industry [1,2]. Efforts have been made in searching for alternative catalysts. Recently, we have successfully acquired Si nanowires with high yield using indium nanoparticles as catalysts by plasma-enhancement [3]. Any potential applications of this type of nanowires into Si-based industry require systematic characterization of structure of the nanowires and deep understanding of their growth mechanisms. Here we report a detailed investigation of morphology and crystalline structure of the In-Si nanowries using high resolution transmission electron microscopy. The results will be compared with the Au-catalyzed nanowires. The Si-nanowires under investigation are synthesized using In nanoparticles with size range of 10 to 200 nm. Figure 1 (a) and (c) show that the droplet at the nanowire tip display spherical shape. The large contact angle to the Si nanowire can be attributed to the lower solubility of Si in indium than in gold. All wires show a base-to-droplet tapering characteristics. It is found that the tapering tendency depends strongly on the size of In droplets: the smaller the droplet, the larger the tapering. Statistical analysis shows that there is an approximate 2:1 ratio of the diameter of In-droplet to the wire diameter at the thin end. The growth orientation of the nanowires is also depending on the size of the droplet. The large droplets prevail in growth along the [111] direction. An example is given in Figure.1 (a) and (b), where the droplet diameter is 105 nm and the wire diameter is 47 nm at the thin end. The growth orientation changes from [111] to [110], as the size of the droplet decreases to around 60 nm. A typical example is given in Figure.1. (c) and (d) for the wire with the droplet size of 62 nm and the thin end wire diameter of 32 nm. The change of Si nanowires growth direction has been reported previously for Au-catalyzed Si nanowires. However, the crossover of growth orientation on Au-catalyzed Si nanowires is smaller, which is usually below 20 nm [4]. The difference may be attributed to the characteristic morphology of Si nanowire grown by using different catalysts. Indium has different surface energy and eutectic compositions as to Au, which may shift the growth-direction transition point. Twins are often observed in the nanowires. The inset in Figure.1 (b) displays the diffraction pattern from {111} twins, showing clearly this nanowire twins on both
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{111} planes (see the markers A and B). The dislocation in the nanowires is not common, except at either the sidewall of the wires or the interface between the droplet and the wire. One of such dislocation is shown in Figure.1 (d) in a twin-free nanowire. In summary, our work suggests that, through fine control of the size of Innanoparticle catalysts, it is possible to control the morphology and growth orientation of the Si nanowires. 1. 2. 3. 4.
J. B. Hannon, S. Kodambaka, F. M. Ross, R. M. Tromp, Nature 440 (2006) 69. J.E. Allen, E.R. Hemesath, D.E. Perea, J. L. Lensch-Falk, Z.Y. Li, F. Yin, M.H. Gass, P. Wang, A.L. Bleloch, R.E. Palmer, L.J. Lauhonn, Nature Nanotechnology, online 10 Feb 2008 F. Iacopi, P.M. Vereecken, M. Schaekers, M. Caymax, N. Moelans, B. Blanpain, O. Richard, C. Detavernier, H. Griffiths, Nanotechnology 18 (2007) 1. Y. Wu, Y. Cui, L. Huynh, C.J. Barrelet, D.C. Bell, C. Lieber, Nano Letters 4 (2004) 433.
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Figure 1. TEM characterization of In-catalyzed Si nanowires. (a) A nanowire with its droplet size of 105 nm. (b) A magnified view of the marked area in (a) and SAED pattern (insert). (c) A nanowire with its droplet diameter of 62 nm. (d) A magnified view of the marked area in (c), SAED pattern (left) and one-dimensional FFT filtered image corresponding to the top right area in (d).
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HRTEM contribution to the study of extraterrestrial nanocarbons and some earth materials analogues J.N. Rouzaud1 and C. Le Guillou1 1. Laboratoire de Géologie de L’Ecole Normale Supérieure, UMR 8538, Paris, France [email protected] Keywords: carbon nanoparticles, extraterrestrial carbon, earth analogues
Most forms of natural and anthropic terrestrial carbons can be found in the extraterrestrial carbons, and especially disordered carbons, graphite, nanodiamonds, onions, … They are often found as nanoparticles and will be named here ‘nanocarbons’. Primitive meteorites contain organic matter of unknown origin (similar to the terrestrial kerogens), and submitted to low temperature metamorphism, but also nanodiamonds and graphitized particles. Meteorites processed at high temperature (1000°C) gather graphitic carbons, stricto sensu graphite, and (nano)diamonds. Thus, meteorites frequently show surprising associations of carbon phases usually synthesized in very different conditions of temperature and pressure [1]. To explain such paradoxical associations, different steps of meteorite processing must be involved (condensation of the protosolar gas, accretion, metamorphism on the parent-body, differentiation, shocks, …) and blends of different components are frequently put forward. To better understand the carbon components formation in Nature, our approach is to compare these natural carbons with possible synthetic earth materials analogues, obtained in well constrained conditions. The aim is to simulate the processes to which carbonaceous matter could have been submitted during their History in space. So, studies were performed on nanocarbons formed by different processes : pyrolysis with or without high pressures, laser pyrolysis, detonation, shocks, chemical vapour deposition, iron catalysed graphitization … The multiscale organisation of these natural and synthetic nanocarbons was studied by High Resolution Transmission Electron Microscopy (HRTEM) with a special interest for the nanometre scale. For disordered carbons, an original image analysis was developed to obtain more quantitative structural data [2]. Some examples of possible earth analogues will be presented: formation of nanodiamond from graphite heated at 1500°C under 15 GPa (Figure 1a from [3], graphitisation of detonation nanodiamonds (Figure 1b). Our approach will be especially illustrated by the comparison between the carbons found in an ureilite (enigmatic thermally processed meteorite) and those obtained by a ‘simple’ shock loading on pure graphite. Ureilites are characterized by a rather paradoxal association of diamond, lonsdaleite, graphite and disordered carbons. Possible synthetic analogues were obtained by 40 GPa, shock loading experiments on graphite [4]. At the nanometre scale, carbon associations are strikingly similar to ureilite carbons. Both natural and analogues samples contain more or less ‘distorted’ graphite grains and nanodiamonds, surrounded by disordered carbon (Figure 2a). The meteorite carbon also shows distorted graphite particles with diamond inclusions (Figure 2b). The
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remarkable parallel contacts between 111 diamond planes and the 002 graphitic ones appear to be typical of a martensittic transformation of graphite into diamond known in shock processes. Such HRTEM observations give constrains on the meteorite history [5]: the diamond of this ureilite was probably formed by a shock on an already graphitised carbon. As far as disordered carbon is concerned, it could be interpreted as a post-shock thermal back-transformation of the newly synthesized nanodiamonds. 1. 2. 3. 4. 5.
Carbon molecules and Materials, éditeurs, ed. R. Setton and P. Bernier, Taylor & Francis, Londres, 2002. A. Galvez, N. Herlin-Boime, C. Reynaud, C. Clinard and J.N. Rouzaud, Carbon 40 (2002), 2775-2789 C. Le Guillou, F. Brunet, T. Irrifune, H. Ohfuji, J.N. Rouzaud, Carbon 45 (2007), 636-648. J.I. Matsuda, A. Kusumi, H. Yajima and Y. Syono, Geochim. Cosmochim. Acta, 59 (1995), 4939–4949] C. Le Guillou, J.N. Rouzaud, L. Remusat, M. Denise and A. Jambon, submitted to Geochim. Cosmochim. Acta et al
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b
Figure 1. HRTEM images of earth analogues; a: 111 planes of diamond formed from graphite (1500°C, 15 Gpa, from [3]), and b: graphitic onion-like surrounding a nanodiamond core (heat-treatment of detonation nanodiamonds at 1100°C)
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Figure 2. HRTEM images of natural carbons from an ureilite meteorite (from [5]); a: nanodiamond (D) covered by disordered graphitic carbon (Dis C); b: HERTEM image and selected area electron diffraction (inset) show the parallelism of the 111 diamond planes (D) and the 002 graphite planes (G).
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Time resolved in-situ TEM observations of Carbon Nanotube growth J. Robertson1, S. Hofmann, R. Sharma2, C. Ducati, R. Dunin-Borkowski 1. Cambridge University, Cambridge CB2 1PZ, UK 2. Arizona State University, USA [email protected] Keywords: environmental TEM, carbon nanotube, catalyst, CVD deposition, growth mechanism
Carbon nanotubes have a unique place in electron microscopy since their discovery by Iijima. They are typically produced by three methods, laser ablation, the arc and by chemical vapour deposition (CVD), usually involving a transition metal catalyst. CVD is most important because it used for industrial scale production, and also on surfaces, for making nanotubes for electronic devices [1]. However, the growth process is not well understood at an atomic level. We have studied the catalyst activation and nucleation of carbon nanotubes using in-situ time resolved transmission electron microscopy, using realistic catalysts and realistic growth conditions and pressures [2]. The catalyst activation is found to consist of a de-wetting of thin layer metal from an oxide support to form nano-crystallites. The catalyst is found to be solid and metallic under our conditions. Nucleation occurs by the formation of a carbon layer cap above the nano-particle which then grows into a nanotube. In CVD typically, one tube grows from each nano-particle. The nanotube walls grow tangentially to the catalyst surface. For multiwall nanotubes, the catalyst particle undergoes severe distortions during grown due to stress, but remains solid. Supplementary in-situ XPS studies confirm that the catalyst is in metallic rather than oxide or carbide state [3]. 1. 2. 3.
J Robertson, Materials Today 10 p36 (Jan 2007) S Hofmann et al, Nanoletts 7 602 (2007) C Mattevi, S Hofmann, et al, J Phys Chem C (accepted)
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Figure 1. The catalyst restructures into a series of nano-particles in order to be active in growth.
Figure 2. In-situ TEM of nucleating and growing carbon nanotubes.
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Insulator-Metal transition: formation of Diamond Nanowires in n-type Conductive UNCD films R. Arenal 1, O. Stephan2, P. Bruno3, D.M. Gruen3 1. LEM, CNRS-ONERA, 92322 Châtillon - France 2. LPS, Université Paris-Sud, 91405 Orsay - France 3. MSD, Argonne National Laboratory, IL 60439 Argonne - USA [email protected] Keywords: Nanodiamond, Nanowires, UNCD, HRTEM, EELS, Transition Insulator-Metal
Under normal process conditions UltraNanoCrystalline Diamond (UNCD) films are highly electrically insulating, but they can become highly conducting when Ar is substituted in the synthesis gas with some of N2. The potential utility of these films that provide the only currently available source of n-type diamond material conducting at ambient temperatures makes interesting to gain a better understanding of the mechanism underlying the insulator-metal transition [1]. In this communication, we present a detailed study of the microstructure of UNCD films as a function of plasma N2 content using electron diffraction, HRTEM and EELS. We performed this work in a FEI Tecnai F20 and in a VG HB501 operating at 200 and 100kV, respectively. The latter microscope provides appropriate conditions to obtain spectroscopy information at the nanometre scale on individual nanostructures [2,3]. Figure 1 (I) and (II) correspond to general views of the samples where elongated nanowires (NWs) are visible. These NWs starts to appear when the N2 content in the gas phase reaches about 10% in volume. They are ~ 100 nm in length and about 5 nm in width as shown in the TEM image (Figure 1 (II)) and more precisely in Figure 2 (I). The lattice fringes corresponding to the d-spacing of the (111) planes of diamond are visible in this latter HRTEM image [1]. The NWs are enveloped by an amorphous layer of about 1nm thick and are embedded by a matrix composed by randomly oriented 3-5 nm crystallites of UNCD. C-K edge is shown in Figure 2 for EEL spectra acquired on different positions on a sample containing those NWs. These spectra were recorded respectively on: the middle of a NW (I), the matrix (II) and the edge of a NW (III). The analysis of the fine structure of the C-K edge confirms that each NW is diamond and that they are enveloped in a sheath of sp2 bonded carbon [1,4]. Moreover we have detected the presence of nitrogen in the amorphous layer. The concentration of N is under one atomic percent. From this study we concluded that the insulator-metal transition of these films is strongly correlated with the formation of these diamond NWs. These NWs are enveloped by an amorphous carbon layer that seems to provide the conductive path for electrons.
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1. 2. 3. 4. 5.
R. Arenal, P. Bruno, D.J. Miller, M. Bleuel, J. Lal, D.M. Gruen, Phys. Rev. B (2007). R.F. Egerton, EELS in the Electron Microscope. 2nd edn. Plenum, New York (1996). R. Arenal, O. Stephan, M. Kociak, D. Taverna, A. Loiseau, C. Colliex, Phys. Rev. Lett. (05). R. Arenal, G. Montagnac, P. Bruno, D.M. Gruen, Phys. Rev. B (2007). This work was supported by the U. S. Department of Energy, Office of Science, under Contract DE-AC02-06CH11357.
Figure 1. (a) SEM and (b) low magnification TEM images of the samples showing the presence of NWs.
Figure 2. (I) HRTEM image showing one of these Diamond NWs and the matrix. (II) C-K edge for several probe positions on an n-UNCD sample corresponding with: (a) middle of a diamond NW, (b) at the edge of the NW and (c) in the matrix, respectively.
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Field emission from iron-filled carbon nanotubes observed in-situ in the scanning electron microscope K.J. Briston1, Y. Peng1,3, N. Grobert2, A.G. Cullis3 and B.J. Inkson1 1. Department of Engineering Materials, University of Sheffield, Mappin Street, Sheffield, UK, S1 3JD 2. Department of Materials, University of Oxford, Parks Road, Oxford, UK, OX1 3PH 3. Department of Electronic and Electrical Engineering, University of Sheffield, Mappin Street, Sheffield, UK, S1 3JD [email protected] Keywords: carbon nanotubes, iron, field emission
Carbon nanotubes (CNTs) have been found to be very good field emitters with low turn-on fields and high current densities [1]. This has made them attractive as possible electron sources in, for example, cathode ray tubes and electron microscopes. Numerous studies have been conducted to date on field emission from standard, unfilled CNTs (e.g. [2-4]), which have linked the field emission properties of CNTs to their structural properties. However, there is very little data available on the field emission properties of filled CNTs, particularly individual filled CNTs. Recently work by Chai et al [5] revealed that arrays of CNTs filled with iron oxide had a lower turn-on field than both arrays of nitrogen-doped CNTs and unfilled CNTs. This raises the possibility that iron-filled CNTs could be similarly successful. In this work, individual iron-filled CNTs were mounted on an SEM nanomanipulator (Kleindiek MM3A) with a sharp electrochemically etched gold tip. A second nanomanipulator with a gold tip was moved to a distance of approximately 300nm-600nm from the freestanding end of the CNT to act as an anode during field emission (see Figure 1). Field emission was obtained from 5 separate iron-filled CNTs. The maximum current that was reached with any of the CNTs was 5.38μA. The turn-on field for the 5 CNTs ranged from 114 to 251 Vμm-1, which is comparable to the turn-on field found for unfilled CNTs using the same experimental set up. The turn-on fields measured are higher than in other works due to the reduced field concentration effect at the very small gap distances used in this work [6]. Two of the CNTs were imaged in a JEOL JEM-2010F transmission electron microscope (TEM) following field emission to view their structures and see how the iron was distributed within them. The structures observed were substantially different: one CNT (CNT-1) was relatively straight and uniform containing numerous sub-5nm iron-rich particles while the other (CNT-2) was curved and contained relatively large iron-rich particles up to ~100nm in diameter that distorted the walls of the CNT. In both cases the iron was distributed asymmetrically along the tubes, being mainly located at the bases of the CNTs. In CNT-1 a high density of iron-rich nanoparticles was located at the very tip of the nanotube (see Figure 2) which had a ~30nm carbon coating. In CNT-
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2, cavities existed near the tip which looked to have once contained iron-rich particles. Investigations are still under way to see if other elements are present in the CNTs after field emission along with the iron. The asymmetric distribution of iron in the CNTs suggests that the iron has moved during field emission. Movement of iron inside CNTs under applied current has been observed previously and is thought to be due to thermally assisted electromigration [7]. It is possible the current passing through the CNT during field emission coupled with the very high temperature produced at the tip (up to 2000K at 1μA [8]) is enough to cause electromigration. The fact that one of the CNTs contains iron-rich nanoparticles is also interesting because no as-produced CNTs containing nanoparticles were seen. It seems that iron-rich nanowires inside the CNT (commonly observed in as-produced CNTs) were broken down into nanoparticles during field emission by the high field and temperature. 1. 2. 3. 4. 5. 6. 7. 8. 9.
Bonard J M, Kind H, Stockli T and Nilsson L A, Solid-State Electronics 45 (2001), p. 893. Kaiser M, Doytcheva M, Verheijen M and de Jonge N, Ultramicroscopy 106 (2006), p. 902. Wang M S, Peng L M, Wang J Y, Jin C H and Chen Q, Journal of Physical Chemistry B 110 (2006), p. 9397. Wang Z L, Gao R P, de Heer W A and Poncharal P, Appl. Phys. Lett. 80 (2002), p. 856. Chai Y, Yu L G, Wang M S, Zhang Q F and Wu J L, Chinese Physics Letters 22 (2005), p. 911. Saito Y, Seko K and Kinoshita J, Diamond and Related Materials 14 (2005), p. 1843. Svensson K, Olin H and Olsson E, Phys. Rev. Lett. 93 (2004), p. 145901. Purcell S T, Vincent P, Journet C and Binh V T, Phys. Rev. Lett. 88 (2002), p. 105502. The authors thank the EPSRC for a PhD studentship and funding GR/S85689/01.
Au
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1μm Figure 1. SEM image showing a single iron-filled CNT attached to the gold tip of a nanomanipulator. The gold tip of a separate nanomanipulator is in position to act as an anode in field emission.
Figure 2. TEM image of the tip of CNT1 showing the high density of iron-rich nanoparticles located there after field emission.
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Templated ordering of fullerenes on nanostructured oxide surfaces D.S. Deak, B.C. Russell, D.T. Newell, K. Porfyrakis, F. Silly and M.R. Castell1 1. Department of Materials, University of Oxford, Parks Road, Oxford, OX1 3PH, U.K. [email protected] Keywords: fullerenes, SrTiO3, scanning tunnelling microscopy
The packing arrangement of molecules on surfaces depends on a variety of factors including molecular shape, molecule - molecule interactions, and molecule - substrate bonding. There are numerous ways of ordering molecules on surfaces. The simplest strategy is to choose a clean surface where molecular adsorption leads to a reconstruction, but this only tends to work for small molecules. To influence larger molecules, such as fullerenes, one needs to create ordered adsorption site arrays that will act as a template. Without a template fullerenes will normally form two dimensional close-packed islands, as shown in Figure 1. My talk will cover our current work on schemes to use nanostructured oxide surfaces as templates to guide the molecular packing of fullerenes into non-close packed arrangements. By varying the surface processing conditions, and hence the stoichiometry, in the surface region of (001) and (111) oriented SrTiO3 crystals we can control which reconstructions are formed [1-5]. Extended annealing in ultra high vacuum (UHV) gives rise to the spontaneous growth of surface nanostructures [6-8]. We use Auger electron spectroscopy to show that the nanostructures are due to local phases caused by an enhanced TiOx concentration in the surface region. We have deposited a variety of empty and endohedral fullerenes (C60, C70, C82, Nd@C82, Er3N@C80) on the nanostructured SrTiO3 (001) surfaces and demonstrate that depending on the substrate conditions the molecules will order into close packed monolayers or lines [9,10]. For a particular nanostructure template we also show open grid array packing [11] as shown in Figure 2. Our latest results on SrTiO3 (111) surfaces demonstrate hexagonal ordering of C60 and C70 molecules. A particular template on this surface lines up the long axes of C70 molecules in one direction. 1. 2. 3. 4. 5. 6. 7. 8. 9.
M.R. Castell, Surf. Sci. 505 (2002), p. 1. K. Johnston, M.R. Castell, A.T. Paxton and M.W. Finnis, Phys. Rev. B 70 (2004) p. 085415. F. Silly, D.T. Newell and M.R. Castell, Surf. Sci. 600 (2006) p. L219. D.T. Newell, A. Harrison, F. Silly and M.R. Castell. Phys. Rev. B 75 (2007) p. 205429. B.C. Russell and M.R. Castell. Phys. Rev. B 75 (2007) p. 155433. M.R. Castell, Surf. Sci. 516 (2002) p. 33. D.S. Deak, F. Silly, D.T. Newell and M.R. Castell, J. Phys. Chem. B 110 (2006) p. 9246. H.L. Marsh, D.S. Deak, F. Silly, A.I. Kirkland and M.R. Castell, Nanotech. 17 (2006) p. 3543. D.S. Deak, F. Silly, K. Porfyrakis and M.R. Castell, Nanotechnology 18 (2007) p. 075301.
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10. D.S. Deak, K. Porfyrakis and M.R. Castell, Chem. Comm. (2007) p. 2941. 11. D.S. Deak, F. Silly, K. Porfyrakis and M.R. Castell, J. Am. Chem. Soc. 128 (2006) p. 13976.
Figure 1. Scanning tunnelling microscope image showing close-packed ordering of endohedral fullerenes Er3N@C80. Image width 55 nm.
Figure 2. Scanning tunnelling microscope image of molecular pairing of Er3N@C80 fullerenes on a nanstructured SrTiO3 (001) surface. Image width 31 nm.
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Carbon nanostructures produced by chlorination of Cr3C2 and Cr(acac)3 A. Gómez-Herrero1, E. Urones-Garrote1, D. Ávila-Brande2, N.A. Katcho2,3, E. Lomba3, A.R. Landa-Cánovas4 and L.C. Otero-Díaz1,2 1. Centro de Microscopia y Citometría, Universidad Complutense de Madrid, Av. Complutense s/n, 28040 Madrid, Spain 2. Dpto. Química Inorgánica, Fac. Ciencias Químicas, Universidad Complutense de Madrid, Av. Complutense s/n, 28040 Madrid, Spain 3. Instituto de Química Física Rocasolano, CSIC, Serrano 119, 28006 Madrid, Spain 4. Instituto de Ciencia de Materiales de Madrid, CSIC, 28049 Madrid, Spain [email protected] Keywords: carbide-derived carbon, disordered carbon.
The preparation of nanocarbons and the study of their promising applications have attracted considerable attention in the last years [1,2]. Carbide-derived carbons (CDC) are a well-known group of nanostructured carbon materials, prepared by direct chlorination of the metal carbides according to the following reaction: MCx(s) + y/2 x C(s) + MCly(g). Previously, we have reported the preparation of CDC from Cl2(g) NbC [3] as well as the results of the direct chlorination of metallocenes: ferrocene [4] and cobaltocene [5]. In this work we are presenting the characterization of nanostructured carbon samples obtained from Cr3C2 and from Cr(acac)3 - acac: acetylacetonate, (C5H7O2)- -, mainly by means of TEM and associated techniques (EDX, EELS), employing a CM 200FEG microscope. The sample prepared from Cr3C2 at 400 ºC (reaction time: 60 minutes) consists of remains of the starting precursor encapsulated in a newly generated disordered carbon coating (see Figure 1a). When the reaction temperature is 900 ºC, SEM micrographs indicate that this CDC sample consists of faceted particles with micrometric size (see Figure 1b). TEM images show that most of them are formed by highly ordered graphitic carbon. Remains of the generated CrCl3 are observed as mono-layers between the graphene sheets, as it is shown in Figure 2a, where CrCl3 is imaged along the [010] direction. Besides, carbon particles with disordered structure are also found in this sample in lower proportion. The sample prepared by direct chlorination of Cr(acac)3 at 900 ºC (reaction time of 30 minutes) consists of irregular particles with porous surface according to SEM micrographs. TEM studies indicate that they present a highly disordered structure, with local stacking of ∼3-5 graphene layers (4-5 nm long), as it is shown in Figure 2b. EDX analyses indicate the presence of traces of Cr and Cl. EELS experiments reveal that the carbon bonding in the highly-ordered CDC particles obtained from Cr2C3 at 900 ºC is ∼ 93% sp2-type, a very close value to the ∼90% proportion found in the sample prepared from Cr(acac)3.
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M. Inagaki, K. Kaneko and T. Nishizawa, Carbon 42 (2004), p. 1401. Y. Gogotsi, “Nanomaterials Handbook”, (CRC Press, Boca Raton) (2006). D. Ávila-Brande, N.A. Katcho, E. Urones-Garrote, A. Gómez-Herrero, A.R. Landa-Cánovas and L.C: Otero-Díaz, Carbon 44 (2006), p. 753. E. Urones-Garrote, D. Ávila-Brande, N.A. Katcho, A. Gómez-Herrero, A.R. Landa-Cánovas and L.C. Otero-Díaz, Carbon 43 (2005), p. 978. E. Urones-Garrote, D. Ávila-Brande, N.A. Katcho, A. Gómez-Herrero, A.R. Landa-Cánovas, E. Lomba and L.C. Otero Díaz, Carbon 45 (2007), p. 1699. We kindly acknowledge the financial support from the project MAT2007-63497
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Figure 1. a) TEM image of a typical particle of the sample prepared from Cr3C2 at 400 ºC; b) SEM micrograph of the CDC sample obtained at 900 ºC.
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Figure 2. a) HRTEM image showing rests of CrCl3 (arrowed) in a CDC particle prepared from Cr3C2 at 900 ºC; b) HRTEM image of a carbon particle obtained from Cr(acac)3 (900 ºC), showing disordered structure with local stacking of graphene layers (arrowed). The corresponding electron diffraction patterns of the whole particle are inset
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Structural peculiarities of carbon onions, formed by different methods B.A. Kulnitskiy, I.A. Perezhogin, V.D. Blank 1. Technological Institute for Superhard and Novel Carbon Materials (TISNCM), Centralnaya Street 7a, 142190 Troitsk, Moscow region, Russia [email protected] Keywords: onion, high pressure, TEM
Carbon onions are concentric multishell fullerene clusters of a few nanometers in size. They can be generated by various techniques. The purpose of the present study is to investigate structural peculiarities of carbon onions, formed under different conditions. The onions in our study were formed by four methods: arc-discharge between two graphite electrodes, high pressure and shear deformation treatment of graphite in the diamond anvil high pressure cell, shock-wave loading of graphite and thermobaric treatment of C60. We examined the onions by using the JEM-2010 highresolution transmission electron microscope, equipped with an EDS detector and the JSM 5800 scanning electron microscope [1, 2]. Following characteristics of carbon onions were examined: shape and amounts of shells, spacing between shells, inner structure and defects. It was shown that mainly high-quality spherical carbon onions, consisting of not more then 10 layers, were produced by the shock-wave loading. Onions, formed by other methods, range from 4 nm to 0.5 µm (5-80 spherical layers). Concentration of defects increased from low value for onions of little size to very high value for onions of big size. The presence of defects is typical for multi-layer onions produced both by high pressure treatment and by arc-discharge. Defects of different types were observed. The amount of concentric shells in onions, formed in diamond anvil high pressure cells, tended to increase with pressure and deformation growth. Large carbon onions, formed under these conditions, consist of more then 60 spherical shells. Applied external pressure resulted in formation of defects in radial direction in onions of big size. Defects represented radial lines, corresponding to the rupture of spherical packing and their displacements by the distance of half interlayer spacing. Ruptured spheres connection can be realized by the formation of sp3-bonds [3]. The spacing between outer shells of the onion as well as the spacing between inner shells is larger than the spacing between intermediate shells. Splitting of some layers in multi-layer carbon onions, produced by arc-discharge, can be explained probably by the non-equilibrium growth process. Very high concentration of defects was found in multi-layer onion-like carbon, formed under thermobaric treatment of C60 (15 GPa, 1100K). 1.
V. D. Blank, V. N. Denisov and A. N. Kirichenko, Nanotechnology, 18 (2007), p. 345601.
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V. D. Blank, A. A. Deribas, and B. A. Kulnitskiy, Materials science Forum 566 (2008), p. 357. A. T. Balaban, D. J. Klein and C. A. Folden, Chem. Phys. Lett 217 (1994), p. 266.
Figure 1. a) Carbon onion, formed by arc discharge, consists of about 80 spherical shells; b) Defects in the spherical packing. c) The diagram, explaining sp3-bonds generation in the case of a “failure” in shell’s periodic arrangement.
Figure 2. a) Carbon onion, formed in graphite at 71 GPa and 5 cycles of shear deformation, consists of 60 spherical shells; b) Radial lines in the onion, shown by arrows, correspond to the rupture of spherical packing and their displacements by the distance of half interlayer spacing. Ruptured spheres connection can be be realized by the formation of sp3-bonds.
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Electron Energy Loss Spectroscopy of La@C82 peapods R.J. Nicholls1, D.A. Eustace2, D. McComb2, G.A.D. Briggs1, D.J.H. Cockayne1 and D.G. Pettifor1 1. Department of Materials, University of Oxford, Parks Road, Oxford OX1 3PH, UK 2. Department of Materials, Imperial College London, London SW7 2AZ, UK [email protected] Keywords: EELS, carbon nanomaterials, La@C82
Carbon nanomaterials, such as fullerenes and nanotubes, have potential applications, such as drug carriers [1] and in quantum nanotechnology [2]. In order to use these materials fully, it is necessary to have a good understanding of their properties and how to alter them. Electron energy loss (EEL) spectroscopy provides a probe of the unoccupied density of states (DOS) of a material. Monochromation of the electron source means that it is possible to get high electron energy resolution, which, along with simulated spectra, can be used to obtain a better understanding of the electronic structure of a material. Previously this combination of experiment and theory has been used to look at fullerene crystals [3], and to determine the effect of changes in bond length on the DOS and EEL spectra [4]. We are now extending this investigation to look at the interaction between a carbon nanotube and fullerenes encapsulated inside it. We are using the newly installed monochromated FEI TITAN at Imperial College London to acquire high energy resolution spectra from La@C82 nanocrystals and single-walled nanotubes filled with La@82. This microscope can be operated at 100 kV with energy resolution ~0.15 eV. Monochromation is essential to this study as we are looking for small differences in the fine structure of the edge. A schematic diagram of La@C82 molecules encapsulated inside a single-walled nanotube is shown in Figure 1. The major isomer of La@C82, which is thought to have C2v symmetry, has been used to make the nanocrystals and to fill the nanotubes. The carbon K-edge from La@C82 is shown in Figure 2a, alongside that from C82. Pristine C82 is thought to have C2 symmetry. The π* peak in the C82 appears narrower than that in the La@C82, which is not surprising as there is less symmetry in the La@C82 case. Figure 2b shows examples of two carbon K-edges from the La@C82 peapod sample. Figure 2b(i) is a spectrum from an area of the sample containing lanthanum and Figure 2b(ii) is from an area which does not contain lanthanum. The area which does not contain lanthanum is likely to be an unfilled nanotube. This is supported by the similarity of the spectrum to that of a single walled nanotube. Apart from the relative heights of the π* and σ* peaks, there are two main differences between the spectra which are a shift in the peak labelled A, and a decrease in the peak labelled B in the case without lanthanum. Our combination of experiment and modelling is the only way to understand the origin of these differences.
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A. Hirsch, Physica Status Solidi B 243 (2006), 3209. J.J. Morton et al., Nature Physics 2 (2006), 40. S.M. Lee et al., Chemical Physics Letters 404 (2005), 206. R.J. Nicholls et al., Micron 37 (2006), 449.
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Figure 1. Schematic drawing of (a) a La@C82 molecule and (b) a single-walled nanotube filled with La@C82 molecules.
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Figure 2. Carbon K-edges from: (a) (i) La@C82, (ii) C82; and (b) a sample of nanotubes containing La@C82, including (i) an area containing La and (ii) an area without La.
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HRTEM studies of Y-junction bamboo-like CN-nanotubes I.A. Perezhogin, B.A. Kulnitskiy, V.D. Blank, D.V. Batov, E.V. Polyakov 1. Technological Institute of Superhard and Novel Carbon Materials, Centralnaya str. 7a, Troitsk, Moscow region, 142190 Russia [email protected] Keywords: Y-junction, carbon nanotube, transmission electron microscopy
Carbon branched structures are of interest because of their novel functionality and possible applications. Seamless junctions between carbon nanotubes (CNT) are obtained due to the topological defects (i. e. pentagons, heptagons, octagons, etc.) in the hexagonal lattice. In order to obtain multiwall Y-junction bamboo-like CNx nanotubes we use a technique of resistive heating of graphite in a water-cooled hot isostatic pressure apparatus (HIP). Argon-nitrogen mixture (with molar ratio 1:2.5) was used as a gas medium. We did not place any catalyst material into the reaction volume on purpose, but we used a mullite wool as one of the heat screen parts, and we believe, that the mullite wool was a source of Mg, Si, Ca, and Al found in catalyst particles of nanotubes. CNx nanotubes were characterized by transmission electron microscope JEM-2010 equipped with the EDS (Energy Dispersive X-Ray Spectroscopy) system. The TEM images have shown that the material consists of many Y-junctions with the main CNT stem and CNT branches. In most cases the angles between the branches are close to 120°. The nanotubes are 80—170 nm in diameter and have either one or several Y-junctions. Y-junction nanotubes both with short and long branches were observed. It is interesting that the short branches as well as those grown at the end of the process have a conic shape. Usually nanotubes are divided into identical compartments, but some of the observed nanotubes have the compartments of larger length. It is seen that all the Y-junction bamboo-like CNx nanotubes contain pear-shaped catalytic particles at the end. It is believed that the growth of branching structures occurs mainly due to the change of both physical and chemical properties of catalyst particles caused by different reasons. Undoubtedly, catalyst nanoparticles are liquid at temperatures above 1400 °C and thus can take any shape during the nanotubes growing process. Inhomogeneous concentration of carbon in the bulk of the catalyst particle and on its surface may result in drastic change of growth conditions at different parts of the catalyst, thus resulting in the growth of branching structures (when the growth stops in certain area and begins at the other place of the catalyst). Figure 1 shows that the temperature change due to, for example, the fluctuation of gas flow in the reactor results in the change of the carbon solution and location of precipitation sites. New deposition site is bordered by black arrows. The nanotube branch growth proceeds in the other direction. The growth in the former direction (figure 1) stops because the direction of the diffusion of the carbon atoms has been changed. It is worth to notice, that the growth of the “stem” of nanotube proceeds by the tip-growth mechanism, which means that a nanotube growth direction is opposite to the direction of carbon diffusion to the precipitation site, and the growth of the “branches” follows the base-growth mechanism which implies the coincidence of S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 179–180, DOI: 10.1007/978-3-540-85226-1_90, © Springer-Verlag Berlin Heidelberg 2008
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the abovementioned directions inside a particle. We believe that CN radicals play a decisive role in growing the bamboo-like nanotubes as they take part in the formation of pentagons needed for the precipitation of curved graphene layers on a catalyst particle surface.
Figure 1. Y-junctions formed in the HIP apparatus.
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EF-TEM observation of biological tissue for risk assessment of fullerene nanoparticles K. Yamamoto1, M. Makino1, E. Kobayashi1 and Y. Morimoto2 1. National Institute of Advanced Industrial Science and Technology (AIST), Higashi, Tsukuba, Japan 2. University of Occupational and Environmental Health, Iseigaoka, Kitakyushu, Japan [email protected] Keywords: fullerene, risk assessment, EF-TEM
Industrial applications of nano-carbon materials such as fullerenes, carbon nanotubes, and nanocrystalline diamonds were reported in many fields recently. The toxicity of these nano-carbon materials for the human was not clear, therefore, the toxicity test and risk assessment were very important. Transmission electron microscope (TEM) was powerful technique to study the nano-world, and used both material and biological research. In the case of the observation for the biological cells, the high contrast observation was important, and then, the cell specimens were usually stained with the heavy elements such as U and Pb to increase the contrast of image. However, it was difficult to observe the stained specimen of biological cells containing carbon nanosized carbon materials, because the contrast from carbon nano-sized particles was weak and below the background of the staining heavy elements. Furthermore, the highresolution observation is necessary for the nanoparticles in the biological cells. In this study, the zero loss imaging of the biological tissue with fullerene nano-sized particles is examined by using an energy-filtering TEM with the objective lens for the high-resolution imaging (Zeiss, EM922HR). The acceleration voltage was 200 kV. The electron spectroscopic zero-loss images, which are filtered at a loss energy of 0 eV with an energy window-width of 20 eV, are formed with both unscattered and elastically scattered electrons. The scattering contrasts, Bragg contrasts, and phase contrasts in the filtered zero-loss images are higher than those in the unfiltered images. The in-vivo test of fullerenes was the intratracheal instillation of fullerenes solution in the rat lung. The test solution was the fullerene nanoparticles in the water with 0.1g/l Tween 80 dispersions. The TEM images of fullerenes in the test solution are shown in Figure 1(a) and (b). The diameter of fullerenes is 20 nm, and the fullerene is crystalline. The mount of fullerenes in the solutions were 100 or 200 μg. The fullerene solution was intratracheally instilled in rat lung. The lung tissues after one week, one month, and three months from the instillation were observed by TEM. The lung tissue was fixed using glutaraldehyde and osmium tetroxide solution, and then dehydrated in ethanol, and embedded in epoxy resin. Ultrathin sections were cut on a diamond knife with microtomy. The staining condition of the tissue specimen was examined. The TEM zero loss image of the alveolar macrophages after one week instillation exposure was shown in Figure 2 (a), and the high resolution image of the black particles indexed by an arrow in Figure 2 (a) was shown in Figure 2 (b). According to the
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electron diffraction analysis of these particles, these are identified as the fullerenes. Fullerenes particles keep the fcc crystal structure. As the diameter of fullerenes in Figure 2 (b) was 20 nm, fullerenes keep the particle size. The most of fullerenes particles are observed in the alveolar macrophages and a part of fullerenes are in the alveolar cells. Fullerenes still remained in the alveolar macrophages and the alveolar cells after one month or 3 months instillation exposure.
Figure 1. Low magnification image of the fullerene particles in the test solutions (a). The diameter of the fullerene particle is 20 nm, and the lattice of (111) is observed (b).
Figure 2. Low magnification image of the alveolar macrophages in the rat lung after one week from the instillation (a). (b) is the high-resolution image of the particles indexed by a white arrow in (a). This work was supported by New Energy and Industrial Technology Development Organization (NEDO) of Japan.
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Looking at the surface of catalysts nanopowders J.C. Hernandez, A.B. Hungria, M. Lopez-Haro, J.A. Perez-Omil, S. Trasobares, S. Bernal, P. Midgley, O. Stephan, J.J. Calvino 1. Dept. de Ciencia de los Materiales e Ingeniería Metalúrgica y Química Inorgánica, Universidad de Cadiz, Campus Rio San Pedro s/n, Puerto Real 11510 (Cadiz), Spain 2. Department of Materials Science and Metallurgy. University of Cambridge, Cambridge, CB2 3QZ, UK 3. Laboratoire de Physique des Solides, Universite Paris Sud, 91495 Orsay, France [email protected] Keywords: Catalysts, surface structure, surface composition
Although the bulk it is usually also involved in supporting catalytic processes, it is clear that surface is the key site of these phenomena. Understanding the complex catalytic action requires detailed information about how synthesis, activation or function itself change structural and compositional features of the surface. Most of the techniques currently available at modern electron microscopes are not surface specific but they provide rich information about this particular location. In this contribution we discuss different cases which illustrate some of the potential capabilities of TEM-STEM in the analysis of the structure and composition of the surface of powdered complex multicomponent catalytic materials based on lanthanide oxides. Thus, the complementary use of STEM-HAADF tomography and atomic imaging by HREM and HAADF has proved essential to understand the changes in the redox behaviour of CexZr1-xO2 mixed oxides [1,2]. These materials, used in the formulation of the catalytic converters of the exhaust gases of gasoline-powered cars, improve their oxygen handling capabilities after specific aging treatments in reducing atmospheres at high temperatures. HREM imaging points to the formation in the bulk of a superstructure as responsible of the improvements of the redox properties. On its hand, HAADF has revealed that this superstructure is related to a disorder-order transition in the cationic sublattice of the oxides. Such structural transition nucleates on the surface and evolves throughout the rest of the bulk giving rise, as revealed by 3D-tomographic reconstruction, to oxide nanocrystallites dominated by {111}-type exposed surfaces. Such surfaces have been proved, by STEM-HAADF images, to be Zr-rich variants of an oxidised pyrochlore-type superstructure of the CexZr1-xO2 oxides. Figure 1. The formation/destruction of these Zr-rich, cation-ordered, {111}-surfaces correlates with the changes in the surface chemistry properties of these oxides [3]. The whole set of data to be presented will demonstrate how a combined use of different (S)TEM techniques provide the necessary information to rationalize a really puzzling chemical problem as it is that of the redox behaviour of the Ce-Zr mixed oxide family. A second example will highlight the contribution of nanoanalysis by STEM-XEDS and STEM-EELS. Results obtained in an in-depth characterization of a series of catalysts based on CexPr1-xO2-δ mixed oxides will be presented. These oxides have been studied as redox components of both Metal/CexPr1-xO2-δ and, more complex, Metal/ CexPr1-xO2-δ /doped-Al2O3 catalysts. Nanoanalytical evidences have been crucial to S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 183–184, DOI: 10.1007/978-3-540-85226-1_92, © Springer-Verlag Berlin Heidelberg 2008
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understand the different chemical processes involved in the synthesis of these metalloaded materials and which strongly affect the surface features of both the original oxides and the metal nanoparticles [4]. Likewise, nanoanalytical tools provide clues about the role of alumina dopants to prevent redox activity decay, due to high temperature aging, in the case of the multicomponent Metal/ CexPr1-xO2-δ /doped-Al2O3 materials. In summary, the results to be presented provide a clear picture of how the appropriate use of a combination of Electron Microscopy techniques becomes a unique tool to reveal the subtle, highly localised, surface modification processes which control the macroscopic performance of powder, nanocrystalline, catalysts. 1. 2. 3. 4. 5.
J.A. Perez-Omil, S. Bernal, J.J Calvino, J.C. Hernandez, C. Mira, M.P. Rodríguez-Luque, R. Erni, N.D. Browning, Chem. Mater., 17(17) (2005), 4282. J.C. Hernandez, A.B. Hungria, J.A. Perez-Omil, S. Trasobares, S. Bernal, P.A Midgley, A. Alavi, J.J. Calvino, J. Phys. Chem. C, 111 (26) (2007), 9001. M.P. Yeste, J.C. Hernandez, S. Bernal, G. Blanco, J.J. Calvino, J.A. Perez-Omil, J.M. Pintado, Chem. Mater., 18(11) (2006), 2750. M.P. Rodriguez-Luque, J.C. Hernandez, M.P. Yeste, S. Bernal, M.A. Cauqui, J.M. Pintado, J.A. Perez-Omil, O. Stephan, J.J. Calvino and S. Trasobares, J. Phys. Chem. C (2007) (in press)† We acknowledge the financial support from Ministry of Education and Science of Spain (MAT2005-00333) and Junta de Andalucia (FQM334, FQM110). Part of the Electron microscopy work was carried out at the Electron Microscopy Division of the Central Services of Science and Technology at the University of Cádiz. The authors acknowledge financial support from the European Union under the Framework 6 program under a contract for an Integrated Infrastructure Initiative. Reference 026019 “ESTEEM”.
Figure 1. (left) 3D reconstruction of a Ce0.68Zr0.32O2 nanocrystallite showing octahedral shapes; (right) HAADF image showing alternation of Ce and Zr rich (111) planes in the Ce0.68Zr0.32O2 oxide with pyrochlore type superstructure. The first plane at the surface corresponds to a Zr-rich variant (lower image intensity in the intensity profile below).
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Gathering structural and analytical information on catalysts at sub-nanometer level with TEM F.J. Cadete Santos Aires and M. Aouine IRCELYON, Institut de Recherches sur la Catalyse et l’Environnement de Lyon. UMR 5256 CNRS/Université Lyon 1. 2, Avenue Albert Einstein. 69626 – Villeurbanne cedex. France [email protected] Keywords: catalysts, TEM
Catalytic reactions are based on the adsorption and possible dissociation of reactants at the surface of catalysts. The prevailing role of the catalyst surface implies the use of very divided materials in order to increase the surface/bulk atom ratio leading to better efficiency and specific properties of the catalysts as well as to atom economy. Catalysts are thus generally formed by nanoparticles (metals, oxides, sulphides, …) supported on low-density materials (which are also more or less divided according to the final application). Most of the characterization techniques used to study these materials are not spatially-resolved and can only give global information on the properties of the catalysts. Since catalysts are composed of nanomaterials it is essential to be able to gather information at sub-nanometer level. Transmission electron microscopy, together with associated analytical techniques, is able to probe the properties of materials at subnanometer and atomic levels and gather localized information on morphology, composition and atomic-structure variations within the catalyst. Such information leads to a good knowledge of the catalyst and is essential to understand its catalytic behavior in relation with its local physico-chemical properties. It is thus in many cases a key and unique technique to elucidate the catalytic behavior of the catalysts (specific activity and selectivity, poisoning and deactivation, …) as well as to give directions for the synthesis of new catalytic materials with specific properties (activity, selectivity, multifunction, resistance to poisoning…). Examples illustrating these aspects [1-5] will be given in this presentation. 1. 2. 3. 4. 5.
F.J. Cadete Santos Aires et al., Proc. EUREM XII (Brno), Vol.II (Physical Sciences), 2000, 395-396. M. Aouine et al. (unpublished). M. Aouine et al., Chem. Commun., 2001, 1180-1181. F.J. Cadete Santos Aires et al., Catal. Today 117 (2006) 518-524. R . Mahfouz et al., Applied Surface Science 254 (2008) 5181-5190.
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Figure 1. (a) co-existence of Pd and PdO nanoparticles supported on α-Si3N4 in a fresh catalyst for the combustion of methane ; (b) PdAu nanoparticle synthesized by laser ablation in vacuum and the corresponding EDX spectrum (d).
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Figure 3. Identification of a new active phase in the selective oxidation of propane towards acrylic acid (MoV0.33Te0.22Nb0.11Ox) : (a) experimental image; (b) model ; (c) simulated image.
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Size effect and influence of nanoparticles thickness on order/disorder phenomena in CoPt nanoparticles D. Alloyeau1,2, C. Ricolleau1, T. Oikawa1,3, C. Langlois1, Y. Le Bouar2 and A. Loiseau2 1.
MPQ, Univ. Paris 7 / CNRS, UMR 7162, Bâtiment Condorcet, 75205 Paris, France 2. LEM, ONERA / CNRS, UMR 104, B.P. 72, 92322 Châtillon, France 3. JEOL (Europe) SAS, 1 Allée de Giverny, 78209 Croissy-sur-Seine, France
[email protected] Keywords: CoPt nanoparticles, Order/disorder phenomena, Electron tomography
The present work is focused on the study of the size effect on the thermodynamical behaviour of CoPt nanoparticules. Equiatomic CoPt bulk present a phase transition at 825°C, between a tetragonal ordered phase (L10) at low temperature and a disordered Face Centered Cubic (FCC) structure at high temperature. Order-disorder phenomena and the order state of CoPt nanoparticles are investigated as a function of the size and shape of the nanoparticles using Transmission Electron Microscopy (TEM). The synthesis of CoPt nanoparticles is performed by alternative deposition of Co and Pt atoms by using the pulsed laser deposition technique (PLD) with a KrF excimer laser (λ=248 nm) in a high vacuum chamber. Nanoparticles are elaborated on thin amorphous carbon films. By means of an in situ heating experiment in a TEM [1] we have determined that the ordering and the growth of the particles occur when the temperature is over 600°C and 700 °C respectively. To observe the size effects on the order-disorder phenomena, disordered particles samples with a mean particle size of 2 and 3 nm have been annealed during 1 hour at 650°C, 700°C and 750°C. The nanoparticles structure as a function of their size has been determined using the STEM / NBD technique that we have recently developed [2]. Consistently with previous works [3], we did not find ordered particles smaller than 3 nm (Figure 1) in the sample annealed at 650°C. This result evidences a decreasing of the order / disorder phase transition temperature from 825°C to at least 650°C for nanoparticles size smaller than 3 nm. Moreover, particles with the same size in the substrate plane are found with different structural state i.e. L10 or FCC (Figure 2). The 3D morphology of the nanoparticles has been also studied by electron tomography experiments (Figure 3). The particles exhibit ellipsoidal shapes and depending on their size and the annealing temperature, the nanoparticles are more or less elongated in the plane of the substrate. Using both electron tomography and HRTEM focal series we determined that disordered nanoparticles have smaller thickness (≤ 2nm) than the ordered nanoparticles (≥ 3 nm). These results show that order-disorder phenomena in CoPt nanoparticles are strongly dependent not only on the nanoparticles size but also on their 3D morphology. 1. 2. 3.
D. Alloyeau et al., Nanotechnology 18 (2007) p. 375301. D. Alloyeau et al., Accepted in Ultramicroscopy (2008). T. Miyazaki et al. Physical Review B 72 (2005), p. 144419.
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Figure 1. STEM BF image of CoPt nanoparticles annealed at 700°C (left). Nano Beam Diffraction pattern of nanoparticles (b) and (c) shown in the BF image (right).
Figure 2. HRTEM images of CoPt nanoparticles annealed over 650°C : (a) disordered particles oriented along the <100> zone axis. Ordered particles oriented along the (b) [001] and (c) [110] zone axes.
Figure 3. (a) Tomogram of CoPt nanoparticles annealed at 750°C. (b) Enlargement of the tomogram showing the slicing of the tomogram in the (x,z) plane (horizontal white line) and the (y,z) plane (vertical white line) allowing the measurement of the nanoparticle thickness.
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In situ L10 ordering of FePt nanoparticles P. Bayle-Guillemaud1, M. Delalande2, V. Monnier3, Y. Samson2 and P. Reiss3 CEA-Grenoble, Institut Nanosciences et Cryogénie, 17, Rue des Martyrs 38054 Grenoble cedex 9, France 1. SP2M/Laboratoire d’Etudes des Matériaux par Microscopie Avancée 2. SP2M/Nanostructures et Magnétisme 3. SPrAM/LEMOH [email protected] Keywords: FePt nanoparticles, L10 ordering, TEM
Assemblies of monodisperse FePt nanoparticles (NPs) are promising candidates for next generation data storage media with recording densities beyond 1 Tbit.inch due to their high uniaxial magnetocrystalline anisotropy, which can reach 107 J.m-3 along the [001] direction of the face-centered tetragonal (fct) L10 ordered phase.2 Triggered by the seminal work of Sun and coworkers [1] the chemical synthesis of FePt NPs has been extensively studied for the last few years. As-synthesized NPs exhibit a chemically disordered fcc structure. Therefore a post-synthetic treatment of the NPs, such as thermal annealing, is necessary to transform the fcc structure into the ordered fct structure which exhibits the magnetic anitotropy . This phase transition is only possible for close to equimolar stoichiometry (Fe50Pt50) NPs. In this work we present results of in situ annealing of two populations of NPs. The NPs are chemically synthesised, more details can be found in [2]. Two samples of NPs (I and II) are investigated in this work with different synthesised parameters. The first sample (I) exhibits NPs of 7nm with a core/shell structure as shown on the EFTEM analysis of Figure 1. The core being Pt rich and the shell being a 1 nm Iron oxide component. The sample II is made of bimodal size distribution: both spherical NPs with about 4 and 14 nm diameter respectively are observed. The TEM experiments were done on a 3010 JEOL microscope equipped with a Gatan Image Filter (GIF). In situ annealing was made using a JEOL specimen holder allowing heating up to 800°C. HREM images have been taken during experiment. STEM/HAADF images have been obtained afterwards on a FEI-Titan microscope equipped with a Cs-probe corrector. For both sample, in situ annealing above 500°C reveals L1O ordering of the NPs. The ordering is firstly detected at the shell of the NPs. For sample I, this could be explained by the diffusion of the Fe species to the core of the NPs before to reach a composition close to Fe50Pt50. Figure 2a shows an HREM pictures with lighter planes (001) with a double fringe spacing with respect to the 002 lattice distance which is due to the L10 ordering. The same process is observed for the two populations of sample I (Figure 2b). These latter NPs have also been observed by Z-contrast (Figure 2c). This Zcontrast imaging reveals more easily the fraction of small and large particles ordered
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after the annealing process. A 600°C annealing temperature is necessary to get full ordered NPs. 1. 2.
Sun, S; Murray, C. B.; Weller, D.; Folks, L.; Moser, A. Science. 2000, 287, 1989-1992 V. Monnier, M. Delalande P. Bayle-Guillemaud, Y Samson and P. Reiss, Small (to be published 2008)
Figure 1. EFTEM images and chemical profile across particles obtained before annealing on Fe(L23), O(K) and Pt(M45) edge of the NPs of sample I showing the coreshell structure and sample II showing homogeneous Fe and Pt distribution.
Figure 2. a and b ) in situ HREM image of the L10 ordering of sample I (a) and II (b). The 001 planes showing the L10 ordering are observed in both case after annealing above 500°C. c) shows a STEM/HAADF image showing alternative bright and darker planes due to Z contrast of the Fe and Pt alternative planes of the L10 chemical ordered structure .
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Characterization of indium doped zinc oxide nanorods H. Burghardt1, H. Schmid1, W. Mader1 1. Rheinische Friedrich-Wilhelms-Universität, Institut für Anorganische Chemie, Römerstr. 164, 53117 Bonn, Germany [email protected] keywords: diffraction contrast, EDX, In-doped zinc oxide nanorod,
Nanorods are a specific group of quasi one-dimensional materials. A wide range of chemical compounds have been synthesized by various growth techniques over the past decade. Zinc oxide (ZnO) is an n-type semiconductor with a band gap of 3.44 eV at 4K and shows interesting electrical and optical properties. The properties of ZnO can be widely varied by doping with different metal oxides. In this contribution ZnO nanorods are doped with indium oxide. The nanorods were grown via the vapour-liquid-solid (VLS) process [1], which is a metal catalyst supported vapour deposition mechanism. Au catalyst particles deposited on fused silica was used as substrate. A mixture of ZnO, In2O3 and graphite powder was used as sources for zinc, indium and oxygen vapour, respectively. The experiments were performed in a furnace with two heating zones and with argon as carrier gas. The substrate temperature was held at 950°C. The nanorods grow in the direction of the ZnO c-axis despite a short kinked segment at the end of the rod. (fig. 1). The nanorods are crystalline and exhibit a characteristic domain structure (fig. 2) with inversion domain boundaries (IDBs) known from bulk In2O3(ZnO)m material [2, 3]. There exist two types of IDBs, parallel to basal planes and parallel to prism planes of ZnO, respectively, which appear in a strictly alternating sequence. The boundaries are best distinguished in dark field images by an abrupt change of diffraction contrast (see fig. 2b). It is known from studies on bulk materials that the spacing of basal plane IDBs determines the local In content in these systems [3]. The density of IDBs in the rod shown in fig. 2 yields an In content of 0.3 at.% (with respect to Zn) under the assumption of one monolayer of In per IDB. Careful chemical analysis of the nanorod segment using EDX in STEM yielded 0.6 at.% In which is twice the In content which would be measured in bulk In2O3(ZnO)m material. The study of this discrepancy is the subject of ongoing research. 1. 2. 3.
R. S. Wagner, W.C. Ellis, Appl. Phys. Lett. 4 (1964), 89-90. A. Loewe, Ph.D. thesis, University of Bonn (2001). C.Li.Y. Bando, M. Nakamura, N. Kimizika, J. Electron Microsc 46 (1997), 119-127
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Figure 1. In-doped ZnO nanorod grown along (0001) on holey carbon film with gold catalyst particle at the end.
Figure 2. In-doped ZnO nanorod with inversion domain microstructure and IDBs marked by arrows. Bright field (a) and corresponding (0002) dark field image.
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Adsorbate-induced restructuring on Pt nanoparticles studied by environmental transmission electron microscopy M. Cabié, S. Giorgio and C.R. Henry CINAM-CNRS, Campus de Luminy, Case 913, 13288 Marseille cedex 9, France [email protected] Keywords: Environmental TEM, Pt catalysts, surface restructuring
The reactivity of metal nanoparticles depends on their shape, their structure and their morphology. Furthermore, the adsorption of reactive species on the surface of these catalysts can modify their morphology. These considerations illustrate the need of in situ characterization in order to better understand the elemental mechanisms of catalytic reactions. The in situ imaging of Pt catalysts deposited on amorphous carbon has been performed under gas pressure with environmental transmission electron microscopy (ETEM). The E-TEM microscope is a standard TEM equipped with an environmental sample holder allowing the insulating of the sample towards the vacuum of the TEM column [1]. Perfectly homogeneous Pt catalysts exhibiting well-defined shape have been chosen to follow surface restructuring induced by gas exposure. These nanoparticles obtained by the decomposition of organometallic compounds present different morphologies with a majority of particles with square outlines as illustrated in Figure 1. According to HRTEM and weak beam dark field (WBDF) imaging, most of the particles are cubes truncated with more or less extended {111} and eventually {110} facets. These particles size is in the range between 10 and 30 nm. Morphological changes have been observed on a same cluster during gas treatments similar to those usually applied to supported metal catalysts, i.e. oxidation-reduction treatments under a pressure of 5 mbar. Under hydrogen, a faceting of the particle is clearly evidenced (see Figure 2). After oxidation of this same particle a continuous rounding of the particle profile is observed: the corners become rounded and the edges roughen. This rounding of the particle profile indicates higher index facets are formed (see Figure 2). Results in agreement with this study were reported for supported Pt catalysts observed ex situ by TEM [2, 3]. 1. 2. 3.
S. Giorgio, S. Sao Joao, S. Nitsche, D. Chaudanson, G. Sitja and C. R. Henry, Ultramicroscopy 106 (2006), p. 503. G. Rupprechter and H. J. Freund, Topics in catalysis 14 (2001), p. 1. A. S. Ramachandran, S. L. Anderson and A. K. Datye, Ultramicroscopy 51 (1993), p. 282.
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40 nm
Figure 1. TEM image showing a collection of Pt nanoparticles composed mainly of truncated cubes.
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EELS in monochromated and Cs probe corrected TEM: a relevant alternative to soft-XAS in synchrotron to investigate nanostructures M. Cheynet1, S. Pokrant2, S. Ersen3. 1. Laboratoire SIMaP 1130 rue de la Piscine BP75 Saint Martin d’Hères 38042 France. 2. Carl Zeiss NTS GmbH, 73447 Oberkochen, Germany. 3. TEM group, Forschungszentrum Caesar, 53175 Bonn, Germany. [email protected] Keywords: Cs probe corrected TEM, Monochromator, core and low loss spectra of TiO2 anatase.
Up to recently, soft-X-Ray Absorption Spectroscopy (SXAS) performed in the context of a 3rd generation Synchrotron Radiation Source (SRS), was the natural experimental way to probe element partial density of empty electronic states. Thanks to the coupling to high resolution monochromator, energy resolution around 0.1 eV is currently achieved in XAS spectra, in an energy range as low as 100-1000 eV. Nowadays, with the development of monochromator and electron probe corrector, 0.1eV energy resolution with beam current ten times higher than in traditional machines can also be obtained in the context of Transmission Electron Microscopes. Thus, EELS in monochromated and Cs probe-corrected FEG-TEM is becoming a reliable alternative to soft-XAS. It is even becoming one of the most relevant tools to investigate nanostructures, because of the sub-nanometer spatial resolution achieved in FEG-TEM. In this work, we report on EELS (core-loss and low-loss) experiments performed using the “CRISP” (i.e. a ZEISS Libra200 Field Emission Gun TEM fitted with a monochromator and a Cs probe corrector, a High Angle Annular Dark Field detector (HAADF) and equipped with a 2nd order corrected Omega filter and a Gatan camera) to investigate TiO2 nano-particles. In this instrumental context, EELS is a simple and fast way to get access to structural information with energy resolution below 0.2 eV and spatial resolution below 1nm. This analytical method works especially well in compounds containing d electrons, as for transition elements based compounds. In these cases, the L2,3-edge is very sharp and strongly influenced by crystal field effects, hence the L2,3-edge ELNES structures are specific for a given crystal structure. In the particular case of TiO2, it is well known that several polymorphic forms exist at ambient conditions. The tetragonal rutile and anatase are the most common phases. As expected, several studies, using soft-XAS in synchrotron radiation [1][2][3] have demonstrated that the Ti-L2,3 edge fine structures are specific for each one of these phases. EELS in conventional TEM have also allowed discriminating these two structures; however the energy resolution is too low to allow resolving all the energy transitions observed in XAS spectra or predicted from calculations. In the first part, we will show the benefits obtained by the coupling of a FEG-TEM to a monochromator and a Cs probe corrector. For this purpose the Zero-Loss peak recorded in vacuum with monochromator “on” and “off” are compared. Then, the Ti-
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L2,3 edge recorded on TiO2 nano-particles with the monochromated (on and off) and Cs probe corrected system are compared to spectra recorded with a “conventional” TEM [4][5] (fig.1). In the second part, the EELS Ti-L2,3 spectrum, recorded on TiO2 nano-particles, using illumination angle α=16mrad, collection angle β= 5mrad, dispersion=0.02eV and probe size of 0.19nm measured in Si, are compared to anatase TiO2 synchrotron soft-XAS spectra published in the literature (figure 2). From the analysis of the full width at half maximum (FWHM) of the Ti-L3 t2g band, we show that the energy resolution achieved in monochromated and Cs probe corrected spectra is as high as in XAS spectra (table 1). In the last part, we present the low-loss spectra recorded on TiO2 nano-particles and compare them to low-loss spectra acquired using a 0.75 eV energy resolution machine and to calculated low-loss spectra [6].
Figure 1. Ti-L2,3 edge recorded from anatase using a Philips CM30 (dotted line) and a Zeiss-CRISP monochromator “On” and “Off” (full line).
Figure 2. Ti-L2,3 edge recorded from anatase: XAS spectrum (dotted line); monochromated and Cs probe corrected EELS spectrum (full line).
Table 1. Comparison of the Full Width at Half Maximum of the Ti-L3 t2g band in different instrumental environment (XAS in Synchrotron or EELS in TEM) Experimental configuration FWHM of the L3 t2g edge CRISP (TEM+mono+ Cs probe)Re : 0.23 eV 0.40 eV this work CM30 GIF 2000 (Re : 0.75 eV) 1.20 eV ref .5 XAS (beam line 8 Berkeley) Re : 0.2 eV 0.39 eV ref.1 VG (Cold FEG-TEM) Re : 0.5 eV 0.72 eV ref.4 XAS (beamline22 MAX LAB Lund Sweden) Re 0.3eV 0.48 eV ref.2 XAS (SRRC-HSGM- Taipei, Taiwan) Re 0.2-03 eV 0.46 eV ref.3 1. 2. 3. 4. 5. 6.
S.O. Kucheyev et al., Phys. Rev. B. 69 (2004) 245102. R. Ruus et al., Sol. Stat. Com. 104, n°4, (1997) 199. Y. Hwu, Y.D Yao, N.F. Chen, C.Y. Tung, H.M. Lin, Nanostructured Mat., 9 (1997) 355. R. Brydson et al, J. Phys. : Conden. Matter, 1 (1989) 797. G. Bertoni et al. Ultramicroscopy 106 (2006) 630. M. Launay, F. Bouchet, P. Moreau, Phys. Rev. B, 69 (2004) 035101.
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Atomic-resolution Electron Microscopy at Ambient Pressure J.F. Creemer1, S. Helveg2, A.M. Molenbroek2, P.M. Sarro1 and H.W. Zandbergen3 1. DIMES-ECTM, Delft University of Technology, P.O. Box 5053, 2600 GB Delft, The Netherlands 2. Haldor Topsøe A/S, Nymøllevej 55, DK-2800 Kgs. Lyngby, Denmark 3. Kavli Institute of Nanoscience, HREM, Delft University of Technology, P.O. Box 5046, 2600 GA Delft, The Netherlands [email protected] Keywords: Environmental TEM, ETEM, microelectromechanical systemts, MEMS, gas-solid interactions, nanocrystals, heterogeneous catalysis
High-resolution TEM (HRTEM) is a powerful technique for atomic-scale imaging of nanomaterials that are kept under high vacuum conditions. However, the technique is also used in in situ studies of gas-solid interactions; an application generally referred to as environmental TEM (ETEM). ETEM at the atomic-scale is very demanding because the gas atoms scatter the electron beam and thereby degrade the resolution. A key requirement for ETEM is therefore that the number of gas atoms along the path of the electron beam is limited. The limitation of the number of gas atoms is usually reached by confining the gas around the solid specimen. The confinement is attained by two types of instruments: differentially pumped vacuum systems and windowed cells [1]. Recently, both instruments have provided images showing atomic lattice fringes with spacing smaller than 0.2-0.3 nm. In these experiments the solid specimens were exposed to gasses at pressures up to about 10 mbar and temperatures up to 900 ºC; see e.g. [2-4]. This atomic-resolution ETEM has led to significant new insights into the mechanisms of gas-solid reactions in a variety of nanostructured materials, such as heterogeneous catalysts [2-4]. It is, however, important that the new insights are applied with caution. The gas pressures of a few millibars are much lower than the ambient pressures of 1 bar and more in which many nanomaterials find technological application. Because the state and properties of nanomaterials can depend strongly on the gas environment it is therefore desirable to have an instrument for atomic-scale reaction imaging at higher pressures for research on applied nanomaterials [5]. Here we report on a novel nanoreactor that enables atomic-resolution ETEM of nanomaterials during exposure to heat and gases at ambient pressure [6]. This is achieved by miniaturizing the gas volume and heater into a microelectromechanical system (MEMS). It fits into the tip of a dedicated sample holder that can be used in a normal CM microscope of Philips/FEI. The performance of the nanoreactor was demonstrated by the in situ studies of a Cu/ZnO catalyst for methanol synthesis. The nanoreactor allowed for the direct observation of nanocrystal growth and mobility on a sub-second time scale during heating to 500 ºC and exposure to 1.2 bar of H2. In the
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same environment, HRTEM images showed atomic lattice fringes in the Cu nanocrystals with spacings that demonstrate a spatial resolution of 0.18 nm. We anticipate that the nanoreactor will generate new insights into a variety functional nanomaterials and the way they interact with ambient environments. 1. 2. 3. 4. 5. 6. 7.
E.P. Bulter, K.F. Hale, Dynamic experiments in the electron microscope, Practical Methods in Electron Microscopy, vol. 9, Nthe Holland, Amsterdam (1981). E.D. Boyes, P.L. Gai, Ultramicroscopy 67 (1997), p. 219 S. Giorgio, S. Sao Joao, S. Nitsche, D. Chaudanson, G Sitja and C.R. Henry, Ultramicroscopy 106 (2006), p. 503. P.L. Hansen, S. Helveg and A.K. Datye, Adv. Catal. 50 (2006), p. 77. N.I. Jaeger, Science 293 (2001), p. 1601. J.F. Creemer, S. Helveg, G.H. Hoveling, S. Ullmann, A.M. Molenbroek, P.M. Sarro and H.W. Zandbergen, accepted to Ultramicroscopy (2008). This work is supported by STW, applied science foundation of NWO and the Ministry of Economic Affairs with financial contributions from FEI Company, Haldor Topsøe A/S, and the European ESTEEM project.
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Development of a system for TEM/STEM investigation of air-sensitive materials: Preliminary results on CeO2 reduction behaviour J.J. Delgado, M. López-Haro, J.D. López-Castro, J.A. Pérez-Omil, S. Trasobares and J.J. Calvino Departamento de Ciencia de Materiales e Ingeniería Metalúrgica y Quimica Inorgánica. Facultad de Ciencias. University of Cadiz. Apdo. 40. 11510 Puerto Real (Cádiz). SPAIN [email protected] Keywords: air sensitive materials, STEM, CeO2 reduction
High Resolution Transmission Electron Microscopy is a powerful tool for advanced nanostructured catalyst characterization [1]. The use of structural as well as analytical techniques (HREM, ED,STEM-HAADF, EELS or EDX) has been crucial to perform a thorough structural, morphological and compositional study of a wide variety of materials at the nanometric scale. However, in the case of catalysts characterization, the study of the chemical environment/sample interaction is conventionally carried out by pre-treating the sample in a conventional reactor, depositing the sample on a grid and finally transferring it into the TEM equipment under air. Some of these materials, like pre-reduced oxides, are oxygen sensitive and the previously mentioned experimental design does not allow to observe the material in its native state. One of the most appealing approaches to solve this problem is the use of dedicated environmental transmission electron microscopes (ETEM) which allow in-situ characterization [2]. Nevertheless, although ETEM allow a live monitoring of the dynamics of gas–solid interactions, the currently available in-situ stages can only operate at pressure values in the millibar range. Thus, a challenging gap to be addressed is to study the gas–solid interactions under more realistic operating conditions. The use of an environmental reaction cell specific for an anaerobic-transfer TEM holder is an alternative approach that allows carrying out the pre-treatment of the sample under more realistic conditions and subsequently transferring it to the TEM, under conditions which prevent any ulterior sample modification. Although the dynamic aspects of the gas-solid interactions are lost in this alternative approach, the characterization of those structural and compositional features which are induced by realistic thermo-chemical treatments, which could be lost by interaction with air components but which are not reversed by evacuation or cooling, can be investigated. Here we present details about the development of a system for TEM/STEM investigation of air-sensitive materials which allow treatments under flow conditions at high pressure as well as the preliminary results we have obtained in the study of CeO2 and Rh/CeO2 reduction/oxidation behaviour. Samples of these two materials were heated at different temperatures (200-800ºC) under vacuum as well as under flowing inert and H2/Ar mixture at atmospheric pressure. After evacuation of the reaction cell,
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the samples were transferred into the TEM under anaerobic conditions to study the structural accommodation of the oxygen deficiency induced by the reduction treatments. We should recall that for CeO2, as well as for some other higher rare earth oxides with fluorite structure, an homologous series of reduced phases, LnnO2n-2m, is known. The members of these series are just superstructures of the fluorite cell characteristic of the dioxide, generated by ordering of oxygen vacancies. In the case of CeO2 experimental evidences from neutron diffraction as well as electron microscopy has allowed detecting the presence of some members of this series in reduced samples [3]. Figure 1 shows some HREM and Digital Diffraction Patterns observed in the Rh/CeO2 material after different pre-treatments. The DDPs contain spots which don’t correspond to the fluorite structure but that can be easily explained by considering the formation of the so-called π (CeO1.875) and β(3) (CeO1.833) phases of the homologous series. It should be pointed out that reduced ceria re-oxidises quite fast and thoroughly even at room temperature. Thus, the identification of patches of these reduced phases demonstrates that the equipment is working properly and that allows creating and preserving the material in a reduced state, being thus suitable to investigate air-sensitive materials. Though this is an aspect that deserves a more detailed experimental work, it is also worth emphasizing that these two phases, π and β(3), have not been detected in reduced samples of CeO2. This experimental fact points out the interest of performing detailed TEM studies with samples prepared under well controlled and more realistic conditions. 1. 2. 3.
P.L. Gai and J.J. Calvino; Annu. Rev. Mater. Res. 35 (2005) 465-504; Boyes, E.D. and P.L. Gai. Ultramicroscopy. 67(1-4): p. 219-232 (1997) C. López-Cartes, J.A. Pérez-Omil, J.J. Calvino, J.M. Pintado, L. Eyring y Z.C. Kang. Ultramicroscopy. 80(1): p. 19-39 (1999)
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Figure 1. High-resolution micrographs of a Rh/CeO2 sample after pre-treatment under H2-Ar (left) and vacuum (right) at 800 ºC and subsequent transfer to the microscope under anaerobic conditions. Digital diffractograms are included.
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Characterization of two new zeolites by combining Electron Microscopy and X-Ray Powder Diffraction analyses E. Di Paola, E. Montanari, S. Zanardi and A. Carati Eni S.p.A., Refining & Marketing Division, San Donato Milanese Research Center, Via F. Maritano 26, I-20097 San Donato Milanese (MI) Italy [email protected] Keywords: zeolite, XRPD, SEM, HRTEM
During exploratory zeolite crystallization experiments performed in our labs, two new materials, named ERS-13 and ERS-14, were synthesized. They were obtained in the frame of a wide activity focused on the use of diquaternary alkylammonium dications (diquats) of general formula [R3N-(CH2)n-NR3]2+ as structure directing agents (SDAs). In particular ERS-13 and ERS-14 were obtained by hydrothermal method using N,N’-tetramethylen-bis-(N-methylpiperidinium) dyhydroxide as SDA. The two products were widely characterized by means of X-Ray Powder Diffraction (XRPD), 29Si Nuclear Magnetic Resonance (NMR) [1], Scanning Electron Microscopy (SEM) [1] and High Resolution Electron Transmission Microscopy (HRTEM). ERS-13 crystallizes as thin plates arranged in form of “desert rose”-like aggregates. The XRPD pattern and the 29Si NMR analysis suggested a partial similarity of ERS-13 with the layered material ITQ-8 [2] and, like this one, it undergoes complete breakdown upon calcination at 550°C. The result of HRTEM characterization is shown in Figure 1; work is in progress in order to determine the crystal structure of this material. ERS-14 showed a XRPD pattern similar to that already reported for the zeolite ITQ10 [1], which was claimed to belong to the zeolite beta family [2,3]. However, a new, weak peak at 2θ = 3.7° (λ = 1.54178 Å), related to a periodicity of 23.6 Å, was found in the XRPD pattern of ERS-14. This peak was not observed in the pattern of ITQ-10 and never reported for the zeolite beta family. In agreement with the XRPD analysis, the Selected Area Electron Diffraction showed a periodicity of 47 Å, which is double with respect to that found by XRPD. On the basis of the HRTEM images, this periodicity was attributed to the formation of the polytype C of zeolite beta [3] between not well defined zones (Figure 2). 1. 2. 3.
The authors thank Elisabetta Previde Massara for SEM characterization and Wallace O’Neil Parker Jr. for NMR analyses. M.J. Diaz-Cabanas, M.A. Camblor, Z. Liu, T. Ohsina and O. Terasaki, J. Mater. Chem. 12 (2002), p. 249. J.M. Newsam, M.M.J. Treacy, W.T. Koetsier and C.B. de Gruyter, Proc. R. Soc. London A 420 (1988), p. 375
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Figure 1. HRTEM micrograph (200 kV) of ERS-13 and Electron Diffraction pattern from selected area.
Figure 2. HRTEM micrograph (200 kV) of an ERS-14 crystal; the periodic formation of the polytype C of zeolite beta between not well defined zones is shown.
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Electron beam-induced effects on copper nanoparticles: coarsening and generation of twins D. Díaz-Droguett1, V. Fuenzalida1 and G. Solórzano2 1. Universidad de Chile, Av. Almirante Blanco Encalada 2008, Santiago, Chile 2. PUC-Rio, Rua Marquês de São Vicente 225, Gávea, Rio de Janeiro, Brazil [email protected] Keywords: nanoparticles, copper, twin
This work reports on the in situ-generation of twins due to the incidence of electron beams on copper nanoparticles when they are investigated under transmission electron microscopy (TEM). Copper nanoparticles were synthesized by means of the inert gas condensation (IGC) method, using hydrogen as carrier gas. The IGC method involves the physical evaporation of a material in the presence of a carrier gas at a pressure around 100 Pa [1]. At this pressure the mean free path is in the submillimeter range and the evaporated material cools down by colliding with the carrier gas molecules in its path, leading to nanoclusters collected onto a cold surface. The size of the nanoclusters depends on the evaporated material, the carrier gas and it pressure; lighter gases lead to the production of smaller particles [1,2]. In this research the IGC method was used to generate copper nanoparticles in hydrogen at pressures between 100 and 1200 Pa and evaporation source temperatures in the range from 1320°C to 1730°C. X-ray powder diffraction (XRD) revealed particles of inhomogeneous composition, with metallic copper, Cu2O and CuO in variable amounts. This is attributed to the oxidation of the copper nanoparticles due to the air exposition after venting the chamber to remove the material from the collector surface. The oxidation state of the surface of the particles was corroborated by X-ray photoelectron spectroscopy (XPS). The morphology and size of the particles was analysed by TEM and their cristallinity by selected area electron diffraction (SAED) patterns. Figure 1 shows a bright field and dark field TEM image with agglomerated copper powder on the grid. At higher magnification faceted particles of about 15 nm were observed. These nanoparticles were not stable under electron irradiation: the thermal stimulus caused the coarsening of the nanoparticles and formation of new ones with twins. Since the particles do not exhibit additional structural changes after longer irradiation, we believe that the twins confer stability to the nanoparticles. The sequential changes at different irradiation times are shown in Figure 2. 1. 2. 3.
C. Granqvist, R. Buhrman, J. Apl. Phy.; vol 47, N.5 (1976). M. Turker, Mat. Sci. and Eng. A 367 (2004) 74-81 This work was partially financed by the Chilean government under grant FONDECYT 1070789. D.Díaz-Droguett acknowledges a doctoral fellowship from CONICYT.
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Figure 1. TEM images showing the agglomerated copper powder on the grid. (a) bright field. (b) dark field.
Figure 2. (a), (b) (c), (d) (e), (f), (g) and (h) are TEM images showing the sequential changes at different irradiation times: coarsening of two copper nanoparticles and formation of twins.
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Role of the catalyst and substrate in nucleation and growth of Single Wall Carbon Nanotubes in HFCVD M.-F. Fiawoo1, N. Brun2, A.-M Bonnot3, O. Stephan2 , J. Thibault4 and A. Loiseau1 1. LEM, ONERA/CNRS UMR104, 92322 Châtillon, France 2. LPS, CNRS UMR 8502, Université Paris Sud, 91405 Orsay, France 3. Institut Louis Néel, CNRS , 25 avenue des Martyrs, 38042 Grenoble, France 4. TECSEN, CNRS UMR622, Université Paul Cézanne Aix-Marseille III, 13397 Marseille, France [email protected] Keywords: carbon nanotubes, nuclei, growth, catalyst
Nowadays many kinds of CVD (Chemical Vapour Deposition) processes manage to produce bundles and isolated SWCNTs (single wall carbon nanotubes) from supported catalytic particles. In order to understand the growth mechanism of these tubes and the role of the catalyst in these processes, as it had already been done for the high temperature syntheses [1, 2], we have undertaken a systematic study by TEM (Transmission Electron Microscopy) of the nanotubes and of their relationships with catalytic particles at different stages of their growth. With this aim, nanotubes were grown by HFCVD (Hot Filament Chemical Vapour Deposition) and the synthesis stopped at different times from 40 s to 50 mn. Catalytic particles were prepared in situ by heating, under a hydrogen atmosphere, a few monolayers of cobalt deposited on silica substrates. These substrates have been specially manufactured, prior to the synthesis, in order to make possible the observation of the nanotubes and particles along the synthesis process (Figure 1) [3]. Our observations demonstrate that, whatever the particle size, ranging from 1.4 to 5 nm, nanotubes nucleate either parallel or perpendicularly to the surface of catalyst particles, with in the latter case, no correlation between the diameters of the tubes and the particles (Figure 1). Surprisingly, tube nuclei are observed all along the synthesis but the perpendicular nucleation mode increases with time. This suggests that corresponding nuclei do not evolve towards nanotubes but are frozen. Thanks to electron diffraction and electron energy loss spectroscopy analyses of the particles, we have shown that metal particles chemically react with silica membranes and progressively transform into a silicide (Figure 2). This reaction is presumed to lead to a deactivation of the catalyst and to the abortion of the nanotube growth. The selective death of a kind of nucleus therefore suggests that different kinetics of growth occur depending on the nucleation mode. This conclusion is supported by the distribution in diameter of the tubes, which is found to be much narrower than that of the nuclei. This lack of correlation suggests that some structural rearrangements take place, before the settlement of the final structure of the tube, which depends on the nucleation mode. As a consequence, different growth kinetics occur which compete with the kinetics of metal – membrane reactions leading to the death of the slowest growing nuclei. Finally our
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results also suggest that helicity of the nanotubes is not settled at the nucleation stage, but later on, during the structural rearrangements of the nuclei. 1. 2. 3.
J. Gavillet et al., P.R.L. 87 (2001), p. 275504-1. A. Loiseau al., C. R. Physique, 4 (2003), p. 975-991. M.-F. Fiawoo et al, JNN submitted (2008)
Figure 1. (a): Suspended SWNTs (b): Isolated carbon nanotube. (c): a nucleus with the graphen sheet parallel to tangent of the cobalt nanoparticle. (d): a nucleus with the graphen sheet perpendicular to the tangent of the cobalt nanoparticle, the arrow is pointing out the end of the nucleus.
Figure 2. The EELS spectra shows the silicon edge showing the onset for cobalt silicide in the dotted square and the silica spectra in black. Bellow, HADF image of a nanotube and cobalt nanoparticles on a silica membrane and chemical maps of Co, Si, O edges and the Si onset.
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PEMFC degradation phenomena studied by electron microscopy L. Guetaz1, B. Vion-Dury1 and S. Escribano1 1. CEA – LITEN, LCPEM, 17 rue des Martyrs, 38048 Grenoble cedex 9, France [email protected] Keywords: PEMFC, degradation, electrocatalyst
Proton-exchange membrane fuel cells (PEMFC) that combine hydrogen and oxygen to produce electricity are attractive zero emission power sources for nomad and stationary applications. The heart of PEMFC is the membrane electrode assembly (MEA). It consists of a proton exchange membrane coated on both side with active layers (AL), where electrochemical reactions take place, and with gas diffusion layers. Fuel cell performance is largely controlled by the AL microstructure composed of nanometer-scale metallic electrocatalysts supported on black carbon surrounding by recast ionomer. Today, significant obstacles to the commercialization of PEMFC result from a large degradation of the performances upon operation. It has been evidence that this performance degradation is due to the degradations of the AL components such as catalyst particle coarsening or carbon support corrosion. However, the relative contribution of each component to PEMFC performance loss and the mechanisms by which the different components are degraded are still largely unknown. For a better understanding, it is therefore necessary to analyse in detail the degradation of each AL constituents. We compare in this study MEA microstructure before and after electrochemical aging. Different SEM and TEM techniques are used for this purpose, depending on the microstructural parameter to characterize: -
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Electrocatalysts: the evolution of the catalyst distribution resulting from coarsening mechanism has been studied using conventional TEM or STEM modes. When the electrocatalysts are bimetallic (Pt-M) nanoparticles, HRTEM, HAADF-HRSTEM and EDX microanalysis have been performed in order to observe the nanoparticle structural and chemical composition changes. Carbon support and re-cast ionomer degradation: The HRTEM images of AL (figure 1a) clearly show the catalyst carbon support surrounding by a few nanometer thick ionomer layer. This thin ionomer layer can also be imaged using HAADF-STEM mode (figure 1b). Porosity: the AL porosity is another crucial parameter that provides gas access to the catalyst site. The 2D porosity can be imaged by FEG-SEM in STEM mode (figure 2). Electron tomography is required for a better knowledge of the porosity in three dimensions.
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Figure 1. Catalyst carbon support surrounding by a few nm thick ionomer layer. a) HRTEM image. b) HAADF-STEM image.
Figure 2. STEM image showing the porosity of the active layer in projection. Electron tomography is required for a better understanding of the porosity in three dimensions.
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TEM investigation of magnetite nanoparticles for biomedical applications S. Gustafsson1, A. Fornara2, F. Ye2, K. Petersson3, C. Johansson3, M. Muhammed2 and E. Olsson1 1. Department of Applied Physics, Chalmers University of Technology, SE-412 96 Göteborg, Sweden 2. Division of Functional Materials, Royal Institute of Technology (KTH), SE-164 40 Kista, Sweden 3. Imego AB, Arvid Hedvalls Backe 4, SE-411 33 Göteborg, Sweden [email protected] Keywords: TEM, magnetite
Magnetic nanoparticles (MNP) are interesting for use in a wide range of biomedical applications including targeted drug delivery, immunoassays, hyperthermia and magnetic resonance imaging (MRI) [1]. The performance of MNP in such applications is dependent on the particle properties which are governed by the particle size, shape and chemistry. This abstract presents a transmission electron microscopy (TEM) study of the nanoscale structure and chemistry of magnetite MNP. The analysis was performed using a Philips CM200 TEM equipped with a field emission gun and operating at 200 kV. Energy dispersive X-ray (EDX) analysis was performed using a Link ISIS system and energy filtered TEM (EFTEM) was performed using a Gatan Imaging Filter (GIF) attached to the microscope. Specimens for TEM were prepared by placing a drop of the ferro fluid on a holey carbon coated copper grid. Figure 1 shows an overview of the magnetic nanoparticles and the corresponding size distribution, as calculated from data from at least 300 particles. The mean particle diameter is 20 nm and the size distribution can be considered as narrow. It can be modelled with a normal distribution. Selected area diffraction in the TEM confirms that the structure of the particles is that of magnetite. EDX analysis also shows that the Fe:O ratio is close to that expected for stoichiometric magnetite. EFTEM imaging of the particles indicates the uniform distribution of iron and oxygen (Figure 2). High resolution TEM (Figure 3) shows that most particles are single crystals. These features are desirable from an application point of view. The particles have sizes just above the superparamagnetic limit which means that they are well suited for specific immunoassay applications [2]. 1. 2.
U. Jeong, X. Teng, Y. Wang, H. Yang and Y. Xia, Advanced Materials 19 (2007), p. 33. A. Prieto Astalan, F. Ahrentorp, C. Johansson, K. Larsson, A. Krozer, Biosensors and Bioelectronics 19 (2004), p. 945.
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Figure 1. (a) TEM micrograph showing the magnetite nanoparticles. (b) The size distribution of the magnetite nanoparticles.
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Figure 2. (a) Bright field image (b) Fe jump ratio image (c) O jump ratio image of magnetite nanoparticles.
Figure 3. High resolution TEM image of a magnetite particle.
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Catalytic soot oxidation studied by Environmental Transmission Electron Microscopy S.B. Simonsen1, S. Dahl1, E. Johnson and S. Helveg1 1. Haldor Topsøe A/S, Nymøllevej 55, DK-2800 Kgs. Lyngby, Denmark 2. Nano-Science Center, Niels Bohr Institute, University of Copenhagen, Denmark [email protected] Keywords: ETEM, catalysis, soot oxidation, diesel engine, emission control
At present, the awareness of soot abatement in the exhaust from diesel engines is increasing due to new environmental legislation for exhaust specifications. An attractive approach for effective soot removal includes the introduction of filters on the dieseldriven vehicles and the functionalization of the filters for catalytic soot oxidation by ceria-based materials. Although it is generally accepted that the redox properties of ceria are crucial to the catalytic effect, the detailed reaction mechanism and the location of the catalytic active sites are still matters of debate. In this contribution we present an Environmental Transmission Electron Microscopy (ETEM) study of ceria-catalyzed soot oxidation related to diesel engine emission control [1]. ETEM has become a powerful tool in heterogeneous catalysis due to its ability to directly monitor catalysts in situ during exposure to reactive gases at elevated temperatures [2]. From time-lapsed ETEM image series of soot particles in contact with a CeO2 catalyst, as illustrated in fig. 1, direct observations at the soot-catalyst interface were obtained during exposure to oxidation conditions and provided mechanistic and kinetic insight into the catalytic oxidation reaction. The results show that the catalytic oxidation reaction involved processes, which were confined to the soot-CeO2 interface region, and that the catalytic reaction surprisingly resulted in motion of soot agglomerates toward the catalyst surface, which acted to re-establish the soot-CeO2 interface in the course of the oxidation process. The observed reaction dynamics was found to be consistent with observations from ex situ oxidation experiments and quantitatively in good agreement with previous kinetic measurements. 1. 2.
S.B. Simonsen, S. Dahl, E. Johnson and S. Helveg, J. Catal. 255 (2008), p. 1. P.L. Hansen, S. Helveg and A.K. Datye, Adv. Catal. 50 (2006), p. 77.
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Figure 1. Time-lapsed ETEM images of soot in contact with a CeO2 catalyst during the exposure to 2mbar O2 at 550ºC. The time interval between the images is ~2min. Scale bar , 90nm. The figure adapted from [1].
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Surface and interface structure of ceria supported ruthenium J.C. Hernandez1, S. Trasobares1, J.M. Gatica1, D.M. Vidal1, M.A. Cauqui1, J.J. Calvino1, A.B. Hungria1 and J.A. Perez-Omil1 1. Dpto. Cienica de los Materiales, Ingenieria Metalurgica y Quimica Inorganica, Universidad de Cadiz, Campus Rio San Pedro s/n, Puerto Real 11510 (Cadiz), Spain [email protected] Keywords: ruthenium catalysts, ceria mixed oxides, electron microscopy
Noble metals highly dispersed on cerium mixed oxide supports Ce(M)O2-x (M=Zr,Tb) are currently being investigated, not only due to their relationship with the so called Three-Ways Catalysts (TWC) but also because of their potential applications in other catalytic processes related to the hydrogen production or the elimination of organic pollutants in residual waters from different industrial processes [1]. Among these supported noble metals catalysts, Ru containing materials seems to be more stable while at the same time very active and selective in comparison with other supported noble metals [2]. This contribution is aimed at investigating the structural features of a new family of ceria-based catalytic materials with Ru as metallic phase. Specifically, the work has been focused on the study of the chemical and nano-structural changes induced on Ru nanoparticles by reduction treatments in hydrogen at increasing temperatures, within the range 623 K - 1173 K. We report detailed information about the nature of the metallic phase (metallic particle morphology and crystallographic features of its exposed surfaces), of the supports as well as of the possible interactions between both components, for a series of Ru nanoparticles supported on different cerium mixed oxides (Ce0.5Zr0.38Tb0.12O2-x, Ce0.62Zr0.38O2 and Ce0.8Tb0.2O2) using high resolution electron microscopy (HREM), high angle annular dark field (HAADF-STEM) imaging, electron energy loss spectroscopy (EELS) and X-ray analysis (XEDS). HREM studies showed that in the whole range of reduction temperatures the Ru nanoparticles were well faceted with morphology close to a truncated hexagonal pyramid as well as epitaxial relationships between metallic nanoparticles and the oxide supports through Ru(002) || Ce(M)O2-x(-111) and Ru[010] || Ce(M)O2-x[110], Figure 1. XEDS and EELS studies confirmed that Ru did not seem to be affected by decoration layers on top of the metal particles or by the formation of alloys or intermetallic phases even after reduction at the highest temperature (1173K), in contrast to that observed for other noble metals like Rh or Pt supported on the same Ce0.8Tb0.2O2 support [3]. A remarkable feature of the Ru catalysts treated in hydrogen at the highest temperature, 1173K, is the presence of small "pedestals" under some metallic particles. The frequency of appearance of these peculiar nanostructures under the metallic nanoparticles increases with the Tb-content of the oxide support (Ce0.8Tb0.2O2 >> 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.
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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 [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.
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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 <112>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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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 [email protected] 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 [email protected] 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 <103> and <112> 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.
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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).
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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).
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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
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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 [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.
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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%
calculado Calc.
200
intensity intensidad
intensity Intensidad
exptal.
150 100
150 100 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 [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.
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a)
b)
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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 [email protected] 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] 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 nm
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 [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 [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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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)<γ(110) [4]. 1. 2. 3. 4.
B. Nikoobakht, M. A. El-Sayed, Chem. Mater. 15 (2003), 1957. Z. L. Wang, M. B. Mohamed, S. Link, M. A. El-Sayed, Surf. Sci. 440 (1999), 809. E. Carbó-Argibay, B. Rodríguez-González, J. Pacifico, I. Pastoriza-Santos, J. Pérez-Juste, L. M. Liz-Marzán, Angew. Chem. Int. Ed. 46 (2007), 8983. L. D. Marks, Rep. Prog. Phys. 57 (1994), 603.
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Figure 1. TEM images of Au nanorods in different growth steps, all particles are oriented along the [001] direction in vertical, as indicated by arrows. The corresponding electron diffraction patterns of the displayed particles, obtained in the [110] zone axis, demonstrate their single-crystalline nature; electron diffraction rotations are compensated. Scale bars: 20 nm.
Figure 2. Structural model of Au nanorods growth steps. a) Original nanorod with octagonal cross-section and {100}, {110}, {111} facets. b) Rod with sharp tips, square cross-section, {110} and {111} facets. c) Growth of a sharpened rod into an octahedral particle by deposition of Au atoms (red spheres) along the {110} facets.
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Reactive Diffusion under Laplace Tension in Spherical Nanostructures C. Ene1, C. Nowak1, and G. Schmitz2 1. Institute of Material Physics, Univ. of Göttingen, Friedrich-Hundt-Platz 1, 37077 Göttingen, Germany 2. Institute of Material Physics, Westf. Wilhelms-Universität of Münster, Wilhelm-Klemm-Str. 10, 48149 Münster, Germany [email protected] Keywords: Laplace tension, core shell nanospheres, reactive diffusion, atom probe tomography
Nanostructures are often distinguished by complex geometries different from a planar layer structure. In order to study the influence of a spherical symmetry, as it is e.g. found in core-shell nano-spheres, thin film Al/Cu/Al and Cu/Al/Cu triple layers are deposited on substrates of 25 nm curvature radius. The thermal reaction is studied by atom probe tomography. Due to the outstanding spatial resolution and the ability to analyse three dimensional structures without projection artefacts, the growth of reaction products at the curved interfaces can be studied even on the length scale of a few nanometers. The experiments demonstrate that the reaction rate depends significantly on the deposition sequence of the metals. This observation is naturally explained by Laplace tension. Tip-shaped substrates have been prepared from tungsten wire by suitable electropolishing and subsequent field evaporation to a curvature radius of 25 to 30 nm at the apex. Triple layer systems Al/Cu/Al and also the inverse stacking sequence Cu/Al/Cu were deposited by ion beam sputtering. Single layer thickness was about 10 nm. Specimens were annealed in a UHV furnace at 110 °C for different times. After annealing, they were analyzed by atom probe tomography (see fig. 1). At both interfaces of the triple layers, the reaction of Cu and Al to a metastable Al2Cu phase is observed, and at both interfaces, these reaction develops according to a parabolic growth law. Therefore, the process is apparently controlled by diffusion across the growing product phase. However, as a most remarkable feature, we note that the width of the products at both Al/Cu interfaces are very different, although an identical reaction proceeds at identical temperature. In the case of the Cu/Al/Cu stacking, the interface close to the tip surface, the outer interface, is that of faster growth (see fig. 2). By contrast, if we chose the inverse Cu/Al/Cu stacking the inner interface is observed to be the broader one. The two growth rates differ by a factor two to three. By a quantitative model, we can show that this growth asymmetry is related to the interfacial tension, the so called Laplace tension induced by the curved interfaces, and its influence on the vacancy equilibrium at both sides of the product layer [1]. Since the diffusional transport of Cu and Al is linked to the directed transport of vacancies, a vacancy gradient necessarily accelerates or decelerates the transport across the product layer and thus the growth of the reaction product.
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C. Ene, C. Nowak, and G. Schmitz, Appl. Phys. Lett. (in press)
Figure 1. (a) Electron micrograph of a layer sample deposited on a tungsten tip. (For clarity, a specimen with enlarged Al layer thickness has been chosen.) (b) Tomographic reconstruction of a typical triple layer. Each detected atom appears in the reconstruction as a grey-coded dot (Cu: light, Al: dark). Inside Cu layers, lattice planes are resolved (inset). deposition sequence 100
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Figure 2. Composition profile of a Cu/Al/Cu stack determined perpendicular to the interfaces after 40 min annealing at 110 °C. Inset shows the result of the opposite stacking sequence Al/Cu/Al.
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Preparation and characterization of palladium nanoparticles with various size distributions M. Slouf1,2, H. Vlkova1, D. Kralova1 1. Institute of Macromolecular Chemistry AS CR, Heyrovskeho namesti 2, 16206 Praha 6, Czech Republic 2. Member of Consortium for Research of Nanostructured and Crosslinked Polymeric Materials (CRNCPM) [email protected] Keywords: colloids, gold and palladium nanoparticles, transmission electron microscopy
Metallic nanoparticles and their colloidal solutions in water (hydrosols) or organic solvents (organosols) have been attracting attention of scientists since 19th century [1]. Hydrosols of many metals, such as Au, Ag, Pt and Pd can be prepared either by chemical syntheses (controlled reduction of water-soluble compounds such as H[AuCl4], ref. [2]) or by suitable physical procedures (laser ablation [3], vacuum sputtering [4], etc.). Nanoparticles are employed in chemistry (catalysis, templates for further synthesis), physics (nanosensors, photophysics), biology (markers in TEM, drug carriers), polymer science (fillers of special polymers, nucleation) and in a number of other areas (medicine, industrial technology). In our previous work we developed a technique of preparation Au nanoparticles with pre-calculated size [5]. Prepared Au hydrosols were very stable and the size of nanoparticles could be varied in the range 5-200 nm. The nanoparticles were used for calibration in quantitative QELS experiments (Quasielastic Light Scattering, ref. [6]) and in laser ablation tests [3]. In the present study we try to introduce a reliable method for preparation of Pd nanoparticles with various sizes. These nanoparticles could be used for the above mentioned applications [3, 5, 6]. Moreover, we plan to employ the Au and Pd nanoparticles for testing of resolution, absorption contrast and chemical contrast in TEM and SEM microscopes equipped with special detectors. Chemical syntheses of Pd hydrosols were based on a well-established method consisting in reduction of PdCl2 with sodium citrate [7]. Our goal was to reproduce the methods described elsewhere and investigate the possibilities of obtaining Pd nanoparticles with different sizes and narrow size distributions. Emphasis was put on simplicity and reproducibility of the preparation techniques. Sizes and purity of the nanoparticles were checked by TEM, SAED and EDX. The first experiments showed that both average sizes and size distributions of Pd nanoparticles were slightly different from the literature [7]. Nevertheless, the sizes of the particles could be slightly varied (Fig. 1abc) although size distributions were somewhat broader (Fig. 1def). 1. 2.
M. Faraday, Philosophical Transactions of the Royal Society of London, 147 (1847), 145181. J. Turkevich, P.C. Stevenson, J. Hillier, Discussions of Faraday Society, 11 (1951), p. 55-75.
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3. 4. 5. 6. 7. 8.
O. Dammer, B. Vlckova, M. Slouf, J. Pfleger. Materials Science and Engineering B - Solid State Materials for Advanced Technology. 140 (2007), p. 138-146. F. Lednicky, E. Pavlova, J. Hromadkova, M. Slouf, R. Masirek. Proceedings of: Mikroskopie 2008. Nove mesto na Morave, 16.-17.2.2006. p. 35 (in Czech). M. Slouf, R. Kuzel, Z. Matej. Z. Kristallogr. Suppl. 23 (2006), 319-324. M. Slouf, J. Plestil, H. Synkova et al., Materials Structure, 12 (2005), p. 82-85. J. Turkevich, G. Kim, Science, 169 (1970), 873-879. Financial support through grants KAN200520704 (Academy of Sciences of the Czech Republic) and GACR 203/07/0717 (Grant Agency of the Czech Republic) is acknowledged.
Figure 1. Three different Pd colloids. TEM micrographs (a, b, c) and size distributions (e, f, g). Standard preparation (b, e). Diluted solutions: 2x (a, d) and 0.5x (c, f).
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Electron microscopy for the characterization of nanoparticles D. Sommer1, U. Golla-Schindler1 1. Institute of Mineralogy, University of Münster, Corrensstr, 24, D-48149 Münster, Germany [email protected] Keywords: nanoparticles, HRTEM, SEM, STEM, EELS
The increasing use of nanoparticles in scientific and industrial applications lets the need to a better understanding into the physical and chemical properties of these particles arise. New generations of electron microscopes like field emission SEMs and energy filtered TEMs open a wide variety of applications. To investigate size, size distribution, morphology, agglomeration and chemical composition in the range of nm and the crystal structure in the range of Ǻ, we use a field emission SEM, the Zeiss 1540 EsB and a conventional JEOL 3010 operating at 297 kV equipped with a LaB6 cathode, post-column Gatan imaging filter and a 1K slow-scan CCD camera. The measurements with the SEM were performed using a special STEM detector, which allows to study conventional TEM samples and to detect the reflected signal above the sample as well as the transmitted signal simultaneously. With TEM, the structure information can be obtained by using diffraction patterns that allow to distinguish between amorphous and crystalline areas of the specimen. In addition, chemical information can be obtained by detection of X-rays with an EDX detector or electron energy loss spectroscopy (EELS) especially for light elements. The analyzed particles are produced from a wide variety of materials, e.g. gold, cobalt ferrites, cobalt, single- and multi-walled carbon nanotubes and CdSe quantum dots. The morphology, size distribution and agglomeration was obtained by taking SEM images in the conventional and transmission mode. The accuracy of the size determination can be significantly improved by using the transmission mode (Figure 1). To yield the morphological properties the nanoparticles were characterized by SEM in the conventional and transmission mode and HRTEM. Figure 2 shows an exemplary SEM image of agglomerated gold nanoparticles with different morphologies and sizes. To study crystallinity, high resolution TEM images were taken that additionally enables the measurement of the lattice spacings (Figure 3a). The chemical composition was determined by EDX and EELS (Figure 3b). 1.
We kindly acknowledge the help of Daniele Bonacchi and Fabio Franchini, who provided most of the nanoparticles and the European Comission for financial support (FP6-2004NMP-TI-4- 032731).
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Figure 1: Cobalt Ferrites a) conventional SEM image recorded with the incolumn SE detector b) image in the transmission mode (STEM) recorded with the ETD-SE detector.
Figure 2: SEM image of agglomerated gold nanoparticles with different morphologies and diagram of the size distribution.
Figure 3: a) HRTEM image of a CdSe Quantum dot. b) Electron energy loss spectroscopy of a cobalt nanoparticle.
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Titanium dioxide nanoparticles prepared from TiOSO4 aqueous solutions J. Šubrt, J. Boháček, N. Murafa, L. Szatmáry Institute of Inorganic Chemistry, 250 68 Řež, Czech Republic [email protected] Keywords: titanium dioxide, photocatalyst, nanoparticles
Titanium dioxide nanoparticles were prepared from hydrated precursors, obtained by precipitation of TiOSO4 aqueous solutions by ammonia or alkali hydroxides under controlled conditions. The precipitation of precursors was followed by annealing of the obtained amorphous or nanocrystalline precipitates. The precipitation procedure has been carried out at constant temperature in range 090 °C and constant pH values 3, 6 or 9 either without additives or using additions of organic modification agents, like glycols, alcohols or amines. The structure and chemical composition of the precipitation products were characterized by XRD, SEM, EDS, HRTEM, SAED and XPS; IR and UV-VIS spectroscopy; and also by thermal analysis. The results demonstrated that precipitates of various crystallinity and different particle size and shape were obtained, depending on the precipitation conditions and modifiers used. The obtained precipitates were subjected to annealing under controlled conditions temperature increase rate, annealing temperature and retention time, atmosphere composition, additives used, etc. The results of XRD measurement carried out at increased temperature showed that the originally amorphous or nanocrystalline precursor when subjected to controlled annealing crystallize gradually and, depending on the annealing conditions, powders with nanocrystals size between 4 – 300 nm can be obtained. Above 800 °C presence of rutile was also observed. The starting gel and the annealing products were studied by electron microscopy (HRTEM and SEM). The results were in agreement with the XRD data and showed continuous growth of the anatase nanoparticles in the products with the heating temperature. The starting gel was formed by clusters of 5-10 nm anatase nanoparticles encapsulated within amorphous material whereas the annealing products were formed by anatase nanoparticles size of which depended on the annealing temperature and time. The anatase nanoparticles were aggregated into clusters of particles containing tens of particular nanocrystals (see Fig. 1). The HRTEM observations showed that the crystallization process can be considered as oriented coalescence of the original nanoparticles into bigger crystals. The conditions of synthesis of the initial hydrated precursor affect significantly course of the processes occurring during the annealing and also the properties of the resulting nanocrystalline TiO2. Depending on the conditions used, powders with required phase composition and crystallite size can be reproducibly synthesized. The obtained materials were tested as photocatalysts using photocatalyzed decomposition of 4-chlorophernol as model reaction. The results showed that some annealing products are
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highly photoactive exceeding markedly the activity of the Degussa P-25 photocatalyst (see Fig. 2).
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Figure 2. Photocatalyzed disappearance of 4-chlorophenol. Sample 1 - starting anatase; sample 2 – 700 °C, 30 min; sample 3 – 700 °C, 8 hrs, sample 4 – 800 °C, 30 min; sample 5 – 800 °C, 8 hrs; sample 6 – 900 °C, 30 min.
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Exploring nanoscale ferroelectricity in isolated and interacting colloidal ferroelectric nanocrystals using electron holography Daniel Szwarcman1, Yossi Lereah2, Gil Markovich1, Martin Linck3 and Hannes Lichte3 1. School of Chemistry, Raymond and Beverly Sackler Faculty of Exact Sciences, Tel-Aviv University, Tel-Aviv, Israel 2. Department of Electrical Engineering-Physical Electronics, The Iby and Aladar Fleischman Faculty of Engineering, Tel-Aviv University, Tel-Aviv, Israel. 3. Triebenberg Laboratory, Institute of Structure Physics, Technische Universität Dresden, Germany [email protected] Keywords: Ferroelectric nanocrystals, Electron holography, Piezoresponse force microscopy.
Whereas ferromagnetic nanoparticles, which form a single ferromagnetic domain, have been heavily studied and precise information exists about their magnetic properties, no such counter-part exists for nanoscale ferroelectric materials. Exchange interactions are responsible for the appearance of ferromagnetism even in small atomic clusters consisting of several Fe, Co or Ni atoms [1]. But in ferroelectric crystalline phases, it is yet to be established, what is the minimal cluster/nanocrystal size and geometry, that could support the ferroelectric crystal phase in the absence of external support such as metal electrodes or electric fields from neighbor particles. It is not even established whether a small single domain isolated nanocrystal could exist as a ferroelectrically aligned structure, because of high depolarizing fields, or that other forms of ferroelectric order should occur. Theory and modeling of such systems have been highly controversial and no detailed experimental data on ferroelectricity in small isolated nanocrystals exists. So far only polycrystalline samples and thin films of ferroelectric materials have been thoroughly studied with standard characterization techniques. Off-axis electron holography is an ideal technique for tackling such problem [2,3]. Electron holography allows the electron phase shift to be recovered. Since this phase shift can be related directly to the electrostatic potential, an electron hologram can be interpreted to provide quantitative information about electromagnetic fields with a spatial resolution approaching the nanometer scale. In this study we investigate the size effect of these isolated ferroelectric nanocrystals by electron holography. Isolated BaTiO3 nanocrystals were prepared by a molten hydroxide synthesis. The ferroelectric properties within the nanocrystals were characterized using electron holography above and below the Curie temperature, and also by Piezoresponse Force Microscopy (PFM), which can show the piezoelectric properties of the ferroelectric nanocrystals. Preliminary results from electron
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holography experiments show (figure 1) different behavior of the polarization within different BaTiO3 nanocrystals with a strong temperature dependence on polarization. 1. 2. 3.
I.M.L Billas, A. Chatelain, W.A de Heer, Science 265 (1994), p.1682. H. Lichte, M. Lehmann, Rep. Prog. Phys 71 (2008), p.016102. H. Lichte, P. Formanek, A. Lenk, M. Linck, C. Matzeck, M. Lehmann and P. Simon, Annu. Rev. Mater. Res., 37 (2007), p.539-588. We kindly acknowledge the help of the European community. This study is supported by EC grant No. 29637.
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STEM investigation on the one-pot synthesis of nanostructured CexZr1-xO2-BaO·nAl2O3 catalytic materials J.C. Hernandez1, J.A. Perez-Omil1, J.J. Calvino1, S. Bernal1, R. di Monte2, S. Desinan2, J. Kašpar2 and S. Trasobares1 1. Dept. de Ciencia de los Materiales e Ingeniería Metalurgica y Química Inorgánica, Universidad de Cadiz, Campus Rio San Pedro s/n, Puerto Real 11510 (Cadiz), Spain. 2. Dipt. di Scienze Chimiche, Università di Trieste, 34127 Trieste, Italy. [email protected] Keywords: STEM, CeZrO2 catalysts, one-pot synthesis
In relation with the applications of CeO2-ZrO2 mixed oxides in catalysis, new synthetic routes aim at preparing materials with improved performances in terms of durability and high temperature textural and chemical resistance [1]. Barium-doped alumina (BDA) as well as Barium hexaaluminates (BHA), BaO·6Al2O3, present high thermal stability, this making them suitable supports of highly dispersed catalytic active phases, e.g. in combustion processes [2]. Thus, these materials seem appropriate to prepare supported CeO2-ZrO2 catalysts with improved high temperature stability. In this contribution we have focused on the characterization by (S)TEM of a series of CexZr1-xO2-BaO·nAl2O3 materials prepared using a novel, one-pot, synthesis route. The goal of the characterization work was to obtain an as exact as possible picture of the distribution of the different components of these nanomaterials, as well as to characterize the chemical and structural features of all the phases involved. In order to evaluate the stability of the newly-prepared catalysts against high temperature aging, the CexZr1-xO2-BaO·nAl2O3 samples were studied after being submitted to calcination treatments at increasing temperatures up to 1200ºC. With respect to the deterioration of the catalytic response of these formulations three aspects become of major concern: (1) the formation of very large CexZr1-xO2 aggregates due to sintering; (2) the formation of large amounts of, low specific surface, BHA at the expense of large specific area BA; (3) The formation of BaO-CexZr1-xO2 or Al2O3- CexZr1-xO2 mixed phases which block the redox activity of the CexZr1-xO2 mixed oxide particles. The CexZr1-xO2-BaO·nAl2O3 samples here investigated were prepared by slow addition of a NH3 solution over a HNO3 solution containing appropriate amounts of Ba(NO3)2, Al(NO3)3, Ce(NO3)3 and Zr(NO3)4. This was followed by a treatment with H2O2 and reflux with 2-propanol for 4-6h. After filtering and washing with 2-propanol, the solid thus obtained was dried at 120ºC and further calcined at 700ºC for 5h. Hereon we will call these samples, fresh samples. HREM images of the fresh samples, Figure 1, confirmed the presence of highly dispersed, nanosized (about 7 nm), CexZr1-xO2 crystallites. In the case of samples with Ce/Zr ratios 0.2/0.8 these corresponded to the tetragonal phase. Alumina is present as γ-Al2O3, with crystal sizes about 6 nm. Increasing the calcination temperature up to 1100ºC, increased the Ce-Zr mixed oxide crystallite size. STEM-XEDS analysis in spot
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mode indicated nevertheless that the CexZr1-xO2 crystallites did not contain Ba after calcination between 700 and 1100ºC, Figure 2. Upon calcination at 1200ºC, HREM images confirmed the formation of needle-like BHA crystals covered by a huge number of Ce-Zr mixed oxide nanocrystals, Figure 3. X-EDS element distribution maps, Figure 3, showed that Ba was preferentially incorporated into the BHA crystals. These results indicate that the synthesis route employed favours the effective separation, self-assembly, of the different components in the final material, what finally allows them expressing their intrinsic functionality. Likewise, (S)TEM appears as an essential tool to reveal the exact distribution as well as the details of the composition and structure of the different components involved in complex catalytic formulations, as those investigated in this contribution. 1. 2. 3. 4.
A. Trovarelli in “Catalysis by Ceria and Related Materials”, ed. Imperial College Press: London, 2002. G. Groppi, M. Bellotto, C. Cristiani, P. Forzatti, P.L. Villa, Appl. Catal. A, 104 (1993), 101. A.J. Zarur and J.Y. Ying, Nature, 403 (2000), 65. This works has received financial support from MEC/FEDER-EU (Proyto MAT200500333), Junta de Andalucia (FQM-110 and FQM-334) and the Ramon y Cajal 2003 Program. [010] (002) (101) (200)
5 nm Figure 1. HREM image of a sample calcined at 700ºC showing nanosized tetragonal-Ce0.2Zr0.8O2 particles.
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Figure 3. STEM HAADF image and XEDS element distribution maps obtained on a sample calcined at 1200ºC
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Enhanced stability against oxidation due to 2D self-organisation of hcp cobalt nanocrystals Isabelle Lisiecki1, S. Turner2, S. Bals2, M.P. Pileni1 and G. Van Tendeloo2 1. Laboratoire LM2N, UMR CNRS 7070, Université P. et M. Curie (Paris VI), B.P. 52, 4 Place Jussieu, F-75231 Paris Cedex 05, France 2. EMAT, University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium [email protected] Keywords: self-assembly, TEM, nanocrystals
Recently, a lot of progress has been made in understanding the changes in physical behaviour of 2D self-assembled nanocrystals and the interactive phenomena due to their long-range organisation. As atoms arranged in a different manner in solids exhibit varying physical properties, physical differences can also be expected when nanocrystals are arranged in a disordered or in a highly ordered structure. For example, the first intrinsic magnetic behaviour in fcc ordered supra-crystals of Co nanoparticles has recently been reported [1,2]. In current work, bright field transmission electron microscopy (TEM), electron diffraction (ED), high-angle annular dark field scanning transmission electron microscopy (HAADF-STEM) and energy-filtered transmission electron microscopy (EFTEM) are used to characterise cobalt nanocrystals with varying mesoscopic ordering before and after an oxidation step. Doing so, we reveal enhanced stability against oxidation due to 2D self-organisation of hexagonally close packed Co nanocrystals. The HAADF-STEM image in figure 1a shows Co nanocrystals on an amorphous carbon substrate prior to oxidation [3]. Using amorphous carbon as a substrate gives rise to both arranged areas and non-arranged/isolated particles [3]. The particles have an average size of 7.5 nm ± 0.4 nm. Electron diffraction performed on arranged and nonarranged/isolated crystals simultaneously (figure 1c) only shows sharp rings corresponding to the hcp Co structure. No oxide-shell crystal structure is detected. The absence of an oxide shell before oxidation is confirmed in figure 1b by line scans taken over the particles indicated in 1a. The particles do not show step-like intensity jumps associated with a dense core – less dense oxide shell structure. After oxidation, core-shell particles have been formed (figure 2). Electron diffraction (figure 2c) now shows an extra ring at 2.46Å, attributed to the cubic structure of CoO [4]. Therefore, the shell that has formed is CoO. Within the hexagonal arrangements the core sizes have shrunk to a size of 5 nm and an oxide shell of 1.5-2 nm has formed. The particles outside the hexagonal arrangement however have a much smaller Co core and thicker oxide shell. In most cases the crystal has even completely oxidised. This is confirmed by the line scans in figure 2b, taken of the particles indicated in 2a. The core-shell boundaries are indicated by dashed lines. The Co core of arranged particle A is larger than the core of particle B. Particle C has totally oxidised and shows a lower intensity associated with less dense CoO.
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To summarise, we show evidence of the enhanced stability against oxidation due to the 2D self-organisation of hcp Co nanoparticles. When the particles are ordered, they favour the formation of core-shell Co-CoO nanostructures, whereas nanocrystals in disordered arrays or isolated on the substrate are found to oxidise completely to CoO.
Figure 1. Co nanocrystals before oxidation: a) HAADF-STEM image showing hexagonally arranged and isolated particles; b) Line scan intensity profiles over the particles indicated in a). No core-shell intensity step is visible; c) SAED pattern of arranged and isolated particles showing only rings corresponding to hcp Co.
Figure 2. Co nanocrystals after oxidation: a) HAADF-STEM image showing arranged particles with core-shell structure and isolated particles with smaller core (larger shell) or totally oxidised particles; b) Line scan intensity profiles over the particles indicated in a). The isolated particles have a smaller core or have been totally oxidised compared to the arranged particles; c) Electron diffraction pattern of arranged and isolated particles together showing hcp Co and cubic CoO rings. 1. 2. 3. 4. 5.
I. Lisiecki, P.A. Albouy, M.P. Pileni, Adv. Mater., 2003, 15, 712 I. Lisiecki, D. Parker, C. Salzemann, M.P. Pileni, Chem. Mat., 2007, 19, 4030. I. Lisiecki, C. Salzemann, D. Parker, P.A. Albouy, M.P. Pileni J. Phys. Chem. C, 2007, 111, 12625. I. Lisiecki, M. Walls, D. Parker, M.P. Pileni Langmuir, 2008, (In press) This work has been performed within the framework of IUAP P5/01 of the Belgian government
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Loaded porous Zn4O(bdc)3 (metal@MOF-5) frameworks characterised by TEM S. Turner1, O.I. Lebedev1, F. Schröder2, R.A. Fischer2 and G. Van Tendeloo1 1. EMAT, University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium 2. Anorganische Chemie II – Organometallics & Materials, Ruhr-Universität Bochum, Universitätsstrasse 150, D-44780, Germany [email protected] Keywords: Metal@MOF-5, TEM, Electron Tomography
In recent years, metal-organic frameworks (MOFs) have received much attention because of their high specific surface areas and pore volumes, applicable in gas storage, catalysis, and photovoltaics [1]. Recently, Zn4O(bdc)3 (MOF-5; bdc = 1,4 benzenedicarboxylate) crystals have been loaded with catalytically active material like Pd, Au, Cu and Ru leading to a heightened catalytic activity in olefin hydrogenolysis and methanol synthesis [2]. These metal@MOF-5 materials can be produced by gasphase loading of metal-carrying precursors into the MOF-5 framework. This loading procedure is known to yield a MOF-5 framework loaded with catalytically active nanoparticles in the range of 1-3 nm [3]. The local distribution of these particles within the MOF-5 framework however remains unclear. Furthermore, MOFs are known to be chemically labile and the loading procedure may affect the framework structure. We show that by minimising the electron dose to avoid beam damage [4], transmission electron microscopy (TEM) can be used to characterise this new family of soft materials on a local scale. Bright field TEM and electron diffraction (ED) are used to provide structural information on the metal@MOF-5 crystals. High-angle annular dark field scanning transmission electron microscopy (HAADF-STEM) imaging aids easy recognition of the particles. The spatial distribution of the metal particles throughout the MOF matrix is visualised using electron tomography. Figure 1a shows a bright field TEM image of Pd@MOF-5. The MOF-5 crystals are clearly faceted; a first indication that the metal organic framework is intact after gasphase loading. The embedded Pd particles (inset) have sizes ranging from 1-5 nm, in good agreement with the metal@MOF-5 model [3]. In the HAADF-STEM image in 1b the Pd particles (bright white contrast) appear to be spread throughout the MOF. The diffraction patterns in 1c and 1d show the cubic structure of the MOF-5 matrix, confirming that the framework remains intact. The faint ring in 1d at 2.25Å corresponds to the crystalline Pd particles. Figure 2a displays a typical Pd@MOF-5 crystal with embedded Pd particles up to 5 nm visible (inset). Some agglomerates are present, possibly indicating a local collapse of the MOF. In 2b, the tomographically reconstructed particles are imaged. Figure 2c is a slice taken through the tomographically reconstructed MOF-5 volume. Electron tomography can clearly distinguish between particles within the MOF and particles on the surface. In the present case, the Pd particles are spread throughout the MOF-5 matrix, and are not preferentially found at the surface of the crystals. S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 275–276, DOI: 10.1007/978-3-540-85226-1_138, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. Palladium@MOF-5: a) Bright field TEM image showing a faceted MOF. Inset: Pd particles in the MOF; b) HAADF-STEM image; c) Diffraction pattern of the cubic MOF-5 framework in [001]* zone axis; d) Diffraction pattern of the framework in [-223]* zone axis and a ring corresponding to the 111 spacing in cubic Pd.
Figure 2. Palladium@MOF-5: a) Bright field TEM image; b) Tomographic reconstruction of the Pd-particles in the MOF-5 matrix; c) Slice taken through the reconstructed volume of the MOF-5 crystal. The metal particles are spread throughout the MOF, although some surface agglomerates are present. 1. 2. 3. 4. 5.
Li, H.; Eddaoudi, M.; O’Keeffe, M; Yaghi, M., Nature 1999, 402, 276-279. Hermes, S et al., Angew. Chem., Int. Ed. 2005, 44, 6237-6241 Schröder F. et al., JACS 2008, (In press) Lebedev, O.I. et al., Chem. Mater. 2005, 17, 6525-6527. The authors acknowledge support from the European Union under the Framework 6 program under a contract from an Integrated Infrastructure Initiative (Reference 026019 ESTEEM)
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Growth behaviour of sub-nm sized focused electron beam induced deposits W.F. van Dorp 1,3, C.W. Hagen1, P.A. Crozier2, P. Kruit1, S. Zalkind3, B. Yakshinskiy3 and T.E. Madey3 1. Delft University of Technology, Lorentzweg 1, 2628 CJ Delft, Netherlands 2. School of Materials, Arizona State University, Tempe, AZ, 85287 3. Laboratory for Surface Modification, Rutgers University, Piscataway, NJ, 08854 [email protected] Keywords: nanometer scale, FEBIP, EBID
Developing techniques for the controlled fabrication of nanostructures is a topic of intense research and is critical to exploit the full potential of nanotechnology. Focused electron beam induced deposition (EBID) is a versatile, direct, and resistless fabrication technique [1, 2]. Precursor molecules, introduced into a vacuum chamber, adsorb on the substrate and are dissociated into fragments under the influence of a focused electron beam. The fragments of the precursor molecules stick to the surface and form a deposit. We performed our EBID experiments in an environmental Scanning Transmission Electron Microscope (STEM) with a beam energy of 200 keV, a 0.3 nm probe, W(CO)6 as a precursor gas and amorphous carbon as a substrate. Typical precursor gas pressures during the deposition were 10-3 Torr, and the sample was held at 150 ˚C to reduce contamination from the microscope. We used the annular dark field (ADF) signal for the imaging of the deposits. By recording the ADF signal during deposit growth, we were able to monitor the growth process in situ. Additionally, atomic force microscope measurements were performed on the deposits to determine the volume of the deposits. We deposit arrays of dots (Fig. 1). The smallest deposits we have been able to fabricate have an average full width at half maximum of 0.7 nm [3]. These are the smallest EBID deposits reported to date and at this small scale, the process is not linear anymore. Although all deposits in the array are created with a fixed dwell time, the amount of mass varies from deposit to deposit. This can be seen in Fig. 1 as a variation in intensity in the ADF image. The distributions of mass become relatively wider as the dwell time decreases and the deposits contain less material. The distributions of mass bear close a close resemblance to Poisson distributions (Fig. 2). This indicates that each of the deposits consists of a discrete number of units. The volume of these units was determined by fitting Poisson distributions to the experimental data and combining this with the AFM measurements and was found to be 0.4 nm3. Our work on EBID continues with a study of the adsorption and dissociation chemistry of CH3-Pt-C5H4CH3 on Au(110). Fig. 3 shows results from temperature programmed desorption (TPD) measurements. The first monolayer in contact with the substrate has a higher desorption temperature and a correspondingly higher adsorption energy than condensed multilayers. We plan to use TPD measurements to determine the cross section for electron beam induced dissociation as a function of electron energy.
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S.J. Randolph, J.D. Fowlkes, P.D. Rack, Crit Rev Solid State Mater Sci 31, (2006) 55 W.F. van Dorp, B. van Someren, C.W. Hagen, P. Kruit, P.A. Crozier, Nano Lett 5, (2005) 1303 W.F. van Dorp, C.W. Hagen, P.A. Crozier, P. Kruit, J Vac Sci Technol B 25 (2007), 2210
Figure 1. ADF image of an array of deposits created with EBID. The dwell time of the e-beam per deposit is identical for all deposits. 140e-9 Torr
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Figure 3. Preliminary results from TPD measurements of CH3-Pt-C5H4CH3 on Au(110). The curves represent the rate of desorption as function of temperature for different gas doses (100s exposure at the indicated pressures). Peaks marking the first monolayer (A) and multilayers (B) are indicated.
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Figure 2. (a) Mass distributions for deposits created with different dwell times. Deposit intensities are measured from ADF images. Negative values are due to the image processing. (b) Another set of data, similar to (a), but with deposit intensities converted to volumes using an AFM measurement. (c) Poisson distributions for expectation values (λ) of 1, 2, 3, 5 and 8. The Poisson distributions are similar to the experimental data.
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Ruthenium deposition on CO2 -treated and untreated carbon black investigated by electron tomography M. Wollgarten, R. Grothausmann, P. Bogdanoff, G. Zehl, I. Dorbandt, S. Fiechter and J. Banhart Helmholtz Centre Berlin for Materials and Energy (formerly Hahn-Meitner-Institute) Glienicker Str. 100, D-14109 Berlin, Germany [email protected] Keywords: Ru catalysts, electron tomography
Ruthenium nano-particles modified by selenium were investigated in the last decade as an alternative to platinum which is used as electro-catalyst in fuel cell cathodes. In order to achieve highly loaded Ru-Se catalysts in a well dispersed and active state, selected parameters of the preparation procedure were optimised, such as the morphology and surface chemistry of the carbon support, the active metal precursor compounds, the amount of the selenium promoter and the process parameters of the thermal treatment during catalyst formation [1]. These investigations revealed, that the characteristics of the carbon blacks used differ significantly and have a severe impact on the distribution and size of the ruthenium particles. The role of the carbon black pretreatment by CO2 on the ruthenium deposition properties is investigated in this contribution. Commercial carbon black (Black Pearls by Cabot and Vulcan XC-72 by Cabot) was used as starting material, which was either used as-received or pre-treated in a CO2 atmosphere at a temperature of 950 °C. Ruthenium loading was done by impregnation with RuCl4 and a subsequent heat treatment at 250°C in a H2/N2 atmosphere. Finally, the carbon-supported ruthenium nanoparticles were modified with selenium by impregnation with SeCl4 and a heat treatment at 850°C. Differences in porosity of the carbon support and the distribution of RuSex particles are investigated by TEM tomography in a Zeiss LIBRA 200 microscope. Reconstructed volume data sets are obtained using the IMOD [2-3] software for reconstruction of tilt series. Figure 1 shows typical examples for the observed morphologies: Vulcan XC-72 carbon black has less porous parts on which large catalyst particles are deposited, whereas Black Pearl carbon black has a fine porosity with small and uniformly distributed Ru-particles. 1. G. Zehl, G. Schmithals, A. Hoell, S. Haas, C. Hartnig, I. Dorbandt, P. Bogdanoff, S. Fiechter, Angewandte Chemie / International Edition 46 (2007) p. 7311-7314 2. J. R. Kremer, D.N. Mastronarde, J.R. McIntosh, J. Struct. Biol. 116 (1996) p. 71-76. 3. D. N. Mastronarde, J. Struct. Biol. 120 (1997) p. 343-352. 4. The work present here is part of the project Application Centre for Tomographic Methods in Material Science and is cofinanced by the European Union and the City State of Berlin.
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(b) Figure 1. Reconstructions of a limited tilt series images obtained by TEM tomography of Ru-Se catalysts (white) on CO2 treated soot (dark grey), (a) Vulcan XC-72 and (b) Black Pearls. Differences in the morphology of both carbon supports as well as in distribution and size of the Ru-particles are clearly visible (grid spacing 20 nm).
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Size-dependent crystallinity and relative orientations of nano-Pt/γ-Al2O3 J.C. Yang1, L. Li1, S. Sanchez2, J.H. Kong2, Q. Wang3, L.L. Wang2, Z. Zhang1, D.D. Johnson2, A.I. Frenkel3, R.J. Nuzzo2 1. Mechanical Eng. & Materials Sci., University of Pittsburgh, 848 Benedum Hall, Pittsburgh, PA, USA 2. Chemistry, Univ. of Illinois at Urbana-Champaign, 505 S. Matthews Ave., Urbana, IL, USA 3. Physics, Yeshiva University, 2495 Amsterdam Ave, Suite 501, New York, NY, USA [email protected] Keywords: Pt, heterogeneous catalysis, nanoparticles, γ-Al2O3
Platinum nanoparticles (NPs) on γ-Al2O3 support is a popular active catalytic system. Recently, in situ X-ray Absorption Spectroscopy (XAS) measurements of this system showed negative thermal expansion, where the nearest neighbor Pt-Pt bond distances decreased with increasing temperature. The unusual phenomena must involve the charge transfer interactions between the Pt particle and the γ-Al2O3 support. Preliminary theoretical simulations revealed that the energetically favorable Pt nanoparticle structures, including their bond-lengths, depend quite sensitively on orientation, surface structure and defects on the γ-Al2O3 surface. For example, O vacancies tend to pin the Pt nanoparticle and significantly change their structure, including a decrease of their bond-lengths with increasing temperatures. Therefore detailed knowledge of the atomic structure of both Pt particles and γ-Al2O3, as well as their structural correlations, is needed. Although high-angle annular dark-field (HAADF) technique is widely used in imaging heterogeneous catalytic materials because of the high contrast of the heavy metal relative to the low-Z support, the information from the support is limited, see Fig. 1. Phase contrast given by highresolution transmission electron microscopy (HRTEM) provides sub-nano information of both the Pt and γ-Al2O3 simultaneously. The samples were prepared by impregnating the Pt2+ precursor, Pt(NH3)4(OH)2⋅H2O, on γ-Al2O3, reducing in H2 gas at 573 K in order to remove the ligands to form metallic nanoparticles. The sizes of Pt particle were controlled by the loading amount, where 1 wt% produced an average Pt sizes of ~1nm and heavy loading of 5 wt% produced sizes of ~2.9 nm. The TEM samples were prepared by spreading a drop of Pt/γ-Al2O3 suspension in ethanol onto an ultra-thin Cgrid, naturally dried. The HRTEM observations were carried out with JEM 2100FEG S/TEM, operated at 200 kV. Fig. 2 is an HRTEM image taken from the 1 nm Pt sample. We noted that the crystallinity of the Pt particles is size-dependent. The 1 nm Pt particles or smaller did not show clear evidence for crystallinity. The lack of uniform bond-lengths and order is supported by XAS and theoretical simulations. The 2.9 nm Pt sample showed fcc structure with a = 0.39 nm, where a few particles contained twin boundaries. The HREM image in Fig. 3 is selected to show orientation correlation of a
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Pt particle with its support. The Pt particle with [110] zone axis located on the γ-Al2O3 with [100] axis, rotated with an angle of ~25.5° between Pt (002) and γ-Al2O3 (004) planes. The mis-orientation of Pt particles on γ-Al2O3 suggests weak interfacial interaction of Pt/γ-Al2O3 system. This research is funded by Department of Energy – Basic Energy Sciences (DE-FG02-03ER15476). The TEM characterizations were carried out at Nanoscale Fabrication and Characterization Facility (NFCF) at the University of Pittsburgh.
Figure 1. HAADF image of Pt particles with an average size of ~2.9 nm.
Figure 2. HREM image of Pt particles and γ-Al2O3. The average size of Pt particles is ~0.9 nm.
Figure 3. HREM image of a Pt [110] particles on γ-Al2O3 [100]. (b) and (c) are FFT from the Pt particle and γ-Al2O3, respectively.
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Formation of nanometer-sized porous GaSb particles by vacancy clustering induced by electronic excitation H. Yasuda1, A. Tanaka1, N. Nitta1, K. Matsumoto1 and H. Mori2 1. Department of Mechanical Engineering, Kobe University, Kobe 657-8501, Japan 2. Research Center for Ultra-High Voltage Electron Microscopy, Osaka University, Suita, Osaka 565-0871, Japan [email protected] Keywords: nanoparticle, structural change, electronic excitation
Porous materials play an important role for developments of luminescent, catalytic, hydrogen storage materials and so on, in which the properties can be greatly improved by a large surface-to-volume ratio. If we can prepare a porous structure not only in the bulk materials but also in the interior of individual nanoparticles, it is expected that the porous structure increases the surface area of the nanoparticles. We found that phase changes take place by atom displacements when GaSb semiconductor compound nanoparticles are excited by electron beam [1,2]. In the present work, the formation of porous semiconductor GaSb compound nanoparticles by clustering of vacancies introduced efficiently by electronic excitation and the formation mechanism has been studied by in situ TEM. Preparation of size-controlled GaSb particles was carried out with the use of a double-source evaporator installed in the specimen chamber of a TEM. Electronic excitation experiments and observations were carried out using the same microscope. The TEM used was Hitachi H-7000 operating at an accelerating voltage of 25 kV. The value of electron flux used for excitations was approximately 1.0x1020 e m-2s-1. The temperature of the supporting films was kept at 430 K. Structural changes associated with excitations were observed in situ. An example of the structural changes in GaSb particles by electronic excitation is shown in Fig. 1. Figures 1(a) and (a') show a BFI of particles with the mean diameter of approximately 20 nm before excitation and the corresponding SAED, respectively. As indexed in Fig. 1(a'), the Debye-Scherrer rings can be consistently indexed as those of GaSb which has the zincblende structure. The same area after excitation for 60 s is shown in Fig. 1(b). In the interior of the particles after the excitation, there appear voids with bright contrast. As seen from a comparison of the magnified images Ia and IIa in (a) with Ib and IIb in (b), the diameter of nanoparticles after the excitation increased up to 15 % compared with those before excitation. In the SAED taken after the excitation as shown in Fig. 1(b'), Debye-Scherrer rings of the zincblende structure are recognized again, but the lattice constant increased up to 1.8 % compared with that before excitation. The same area after excitation for 480 s is shown in Fig. 1(c). The voids in the individual particles change in the shape and size, as seen from a comparison of the magnified images Ib and IIb in (b) with Ic and IIc in (c). In the SAED taken after the excitation as shown in Fig. 1(c'), Debye-Scherrer rings are recognized, superimposed on
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a weak halo ring. The Debye-Scherrer rings can be indexed consistently as those of crystalline antimony, which has the hexagonal structure. The value of the scattering vector (K= (4πsinθ)/λ) for the halo ring is approximately 31.0 nm-1 which is corresponding to the first halo of liquid gallium. This result indicates that a two-phase mixture consisting of a crystalline antimony core and a liquid gallium shell was formed in the particles. From these results, it has been evident that when GaSb particles kept at 430 K are excited by 25keV electrons, two-phase separation takes place via void formation. In this case, it is considered that gallium atoms on the lattice points are displaced by the electronic excitation to form vacancies and gallium interstitials in the crystal. Fig. 2 shows schematic illustrations of defect migrations and defect concentrations as a function of distance from the center of a nanoparticle and the structural changes. The vacancy concentration in the particle core is higher than that in the surface layer, but interstitial concentration increases toward the surface. Consequently, under the condition of vacancy supersaturation in the particle core the vacancy clusters will grow to form a void, and the subsequent surface segregation of interstitial clusters will bring about the separation to the two-phase structure. 1. 2.
H.Yasuda, H.Mori, J. G. Lee, Phys. Rev. Lett. 92 (2004), p.135501. H. Yasuda, H. Mori, J. G. Lee, Phys. Rev. B 70 (2004), p.214105.
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Figure 1. An example of the structural changes in GaSb particles kept at 430 K by electronic excitation. (a)(a') before excitation, (b)(b') after 60 s, (c)(c’) after 480 s. Vacancy Interstitial
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Figure 2. Schematic illustrations of defect migrations and defect concentrations as a function of distance from the center of a nanoparticle and the behaviors of the structural changes.
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Structural investigations of membrane electrode assemblies in fuel cells via environmental scanning electron microscopy S. Zils1, N. Benker2 and C. Roth1 1. Institute for Materials Science, Technische Universität Darmstadt, Petersenstr. 23, 64287 Darmstadt, Germany 2. Institute of Applied Geosciences, Technische Universität Darmstadt, Schnittspahnstraße 9, 64287 Darmstadt, Germany [email protected] Keywords: ESEM, PEMFC, membrane electrode assembly, degradation, quasi in-situ structure research
The ideal membrane electrode assembly (MEA) for polymer electrolyte membrane fuel cells (PEMFC) would combine an improved accessibility of the noble metal catalysts with a good electronic contact of the ion- and electron-conducting parts of the fuel cell. To investigate the contact between these ion- and electron-conducting parts, so far detailed studies have only been performed by using Scanning Electron Microscopy (SEM) and Transmission Electron Microscopy (TEM) [1], which exclude the investigation of the MEA under realistic conditions. In this work, environmental scanning electron microscopy (ESEM) is used in order to obtain significant information of the structure under quasi in-situ conditions. This technique allows for studying the membrane electrode assemblies under different gas atmospheres as well as in different relative humidities and temperatures. These aspects render the possibility to adjust conditions in the sample chamber similar to a working fuel cell. Membrane electrode assemblies with a noble metal loading of 1.6 mg cm-2 have been manufactured by using an airbrush technique first described by Wilson and Gottesfeld [2]. Nafion® 117 (DuPont) has been applied as proton conducting membrane. On both the anode and the cathode side carbon-supported platinum catalysts (20 wt% on carbon black) have been used as electrodes. Sample preparation for ESEM was performed by cutting slices with a size of 2 mm by 10 mm out of the MEA. The ESEM measurements were carried out with a Quanta 200 F (FEI company, Netherlands) equipped with a field emission gun (FEG) and an energy dispersive X-ray (EDX) detector for elemental analysis. The pressure in the sample chamber can be increased in steps of 0.01 mbar, while the temperature of the sample can be varied in steps of 0.1 K using a Peltier element. For the quasi in-situ experiments a temperature of 5 °C was maintained. The pressure in the chamber and therewith the relative humidity (RH) were increased stepwise from 0.5 mbar corresponding to a RH of about 5 % up to a pressure of 8.7 mbar and a RH of 100 % at 5 °C, respectively. Figure 1 shows a sequence of two ESEM micrographs of a MEA cross section. The left image shows the cross section of the MEA at a relatively low pressure of 0.5 mbar.
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From this starting point the pressure in the sample chamber was increased stepwise and a swelling of the membrane electrode assembly could be discovered. First droplets on the membrane surface were observed at a RH of about 20 %, first droplets on the electrode surfaces at a RH of about 35 %.
Figure 1: Secondary electron images of a MEA cross section at 0.5 mbar (~ 5 % RH) and at 3.11 mbar (~ 36 % RH). The MEA swelled due to the chemical properties of the perfluorinated polysulfonic acid Nafion® until the maximum content of water in the membrane was reached. The droplets on the surface of the electrodes started to grow only after this point was reached. We expect that the comparison of a reference MEA with a MEA, which has been running in a long-term single cell test, will give us valuable insight into the interface between the electrodes and the proton-conducting membrane as well as about the effect of water at this interface. At long sight, we might even be able to investigate the cold start behaviour of polymer electrolyte fuel cells quasi in-situ and achieve significant information for an optimized fabrication of membrane electrode assemblies. 1. 2.
F. Scheiba, N. Benker, U. Kunz, C. Roth and H. Fuess, J. Power Sources 177 (2008), p. 273. M. S. Wilson and S. Gottesfeld, J. Electrochem. Soc. 139 (1992), p. 1.
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In situ TEM nanocompression testing A.M. Minor1, J. Ye1 and R.K. Mishra2 1. National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, Berkeley, CA USA 2. General Motors Research and Development Center, Warren, MI USA [email protected] Keywords: in situ TEM, mechanical properties
Recent progress in small-scale mechanical testing methods has greatly improved our understanding of size effects in volumes of a few nanometers to a few microns [1-5]. Besides the important results related to the effect of size on the strength of small structures, the ability to systematically measure the mechanical properties of small volumes of material allows us to test samples that cannot easily be processed in bulk form. Recently, we have used a quantitative in situ TEM mechanical testing device to compress microfabricated pillars inside a TEM so that we can make a direct correlation between an imposed stress and an individual deformation event. This talk will demonstrate how individual microstructural features can be tested directly with this technique, allowing us to explore the fundamental origins of strength and ductility through dynamic observations of plasticity in metallic systems (see Figures 1 and 2). Specifically, results will be presented where the in situ experiments give us some insight into the size effects seen in ex situ pillar compression tests. In addition to the studies of size effects in nanoscale volumes, the in situ nanocompression technique has allowed us to probe the fundamental origins of ductility in complex metallic alloys. As an example of this latter capability, two AA6063 alloys with identical compositions were studied for comparison, one as-extruded and the other heat treated to increase the solute concentration in the matrix. By performing in-situ nanocompression tests on microfabricated pillars we found that the annealed sample exhibited a higher yield stress and greater nanoscale ductility as compared to the as-extruded sample, analogous to what is seen through bulk mechanical testing. Furthermore, our dynamic observations and comparison of the two alloys demonstrate that the dislocation plasticity behavior is quite different for the two samples. The annealed sample showed a complex three-dimensional deformation behavior due to extreme dislocation entanglement, whereas the as-extruded sample showed a simpler two-dimensional plasticity composed of mostly planar slip behavior. This difference in behavior was attributed to the lowered stacking fault energy and thus decreased cross-slip in the annealed alloy. Our observations suggest that achieving a more complex threedimensional plasticity at the nanoscale can be directly correlated to increasing the ductility in a bulk alloy. Demonstrating the power of this new in situ mechanical testing capability, we have successfully correlated the bulk mechanical behavior of a complex alloy with its nanoscale plasticity.
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Uchic, M.D., Dimiduk, D.M., Florando, J.N. & Nix, W.D. Sample dimensions influence strength and crystal plasticity. Science 305, (2004), p. 986-989. Greer, J.R., Oliver, W.C. & Nix, W.D. Size dependence of mechanical properties of gold at the micron scale in the absence of strain gradients. Acta Mater 53, 1821-1830 (2005).F. Author, S. Author and T.H. Author, Journal volume (year), p. 1. Volkert, C.A. & Lilleodden, E.T. Size effects in the deformation of sub-micron Au columns. Philos Mag 86, 5567-5579 (2006). A.M. Minor, Z.W. Shan, E.A. Stach, S.A. Syed Asif, E. Cyrankowski, T.J. Wyrobek, and O.L. Warren, "A new view of the onset of plasticity during the nanoindentation of aluminum", Nature Materials, 5, (697-702) 2006 Z.W. Shan, R.K. Mishra, S.A. Syed Asif, O.L. Warren and A.M. Minor, “Mechanical annealing and source-limited deformation in submicron-diameter Ni crystals”, Nature Materials, 7 (115-119) 2007 This research was supported by the Scientific User Facilities Division of the Office of Basic Energy Sciences, U.S. Department of Energy under Contract # DE-AC02-05CH11231 and the General Motors Research and Development Center.
Figure 1. Low magnification TEM image showing in situ nano-compression experimental setup. The flat diamond punch can be seen in the upper right, approaching the top of the pillar structure on the left.
Figure 2. In situ nanocompression test of a AA6063 alloy micropillar. (a) The pillar before compression, (b) stress vs. displacement data from the in situ test, (c) the pillar after compression.
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Physical measurements on an individual nanostructure in a TEM nanolaboratory M. Kobylko1, S. Mazzucco1, R. Bernard1, M. Kociak1, C. Colliex1 1. Laboratoire de Physique des Solides, UMR CNRS 8502, Bât. 510, Université Paris-Sud 11, F-91405 Orsay, France [email protected] Keywords: in situ microscopy, transport measurements, dedicated specimen holders
With the rapidly growing development in nanosciences, the requirement for realizing physical measurements on well-identified and -characterized individual nanostructures has become unavoidable to complement the vision brought by average measurements on an ensemble of insufficiently characterized nanostructures. A TEM column offers the right environment to explore this approach. It is as more of present concern as two factors favour this technical development: (i) the limited available space between the objective lens pole piece gap has been relaxed with the introduction of Cs correctors in the TEM column (as an example a probe size diameter of 0.1nm is achieved in a SuperSTEM Nion column with a 4mm gap); (ii) advances in microtechnology has made the miniaturization of actuators and sensors more widely available. Within this global context, we have become fully convinced that the design and fabrication of in situ devices is a priority for preparing the way towards a more comprehensive use of future TEM. These instruments should offer all components of a nanolaboratory, where the experimentalist can chose a nanostructure of given size, dimension, structure and composition as the object to be studied during application of an external constraint such as an electric, a magnetic or an electromagnetic field. In this contribution, our intent is to illustrate how we approach this goal by designing and using in situ transport and nanomanipulation TEM sample holders, following some of the trends already explored by Olin and co-workers [1]. The first one is targeted for simultaneous observation / structure determination and electrical transport measurements on on chip nanoobjects (Figure 1a). The second one allows simultaneous observation / structure determination, electrical transport measurements and mechanical manipulation of nanoobjects placed on tips (Figure 1b). Other devices have been built for in situ detection (or illumination) of light from or on a given area of a nanostructure [2]. These in situ TEM-holders pave the way to new meaningful experiments. For example, using a home-made TEM holder of type 2, we could monitor the process of dipping individual CNTs into mercury droplets. Such experiments are intended to measure the transport properties of individual CNTs (ballistic or diffusive behaviour, conductance quantification, etc). Our in situ experiments show that the measurement of conductance quantification in CNTs can be related to the formation of Hg-nanocontacts (Figure 2) and not to an intrinsic property of the CNTs, as suggested by earlier studies [3]. In addition, elastic medium theory simulations reveal that, when CNTs of small radii or made of few walls are considered, the probability that the CNT
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actually enters the drop of mercury is very small [4]. These are only first results confirming the importance of developing the “nanolaboratory” approach within a TEM column, but there is no doubt that many others will rapidly follow. 1. 2. 3. 4.
K. Svensson, et al., Review of Scientific Instruments 74 (2003), 4945. See Mazzucco et al, these conference proceedings. S. Frank, P. Poncharal, Z.L. Wang and W.A. de Heer, Science 280 (1998), 1744. M. Kobylko, M. Kociak, et al., to be published (2008).
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(b)
Figure 1. TEM-holders for in situ TEM measurements: (a) Transport holder for simultaneous electrical transport measurements and structure determination/observation of on chip nanoobjects; (b) Nanomanipulator holder for simultaneous electrical transport measurements, structure determination/observation and mechanical manipulation of nanoobjects on tips.
Figure 2. In situ measurement of the conductance between a tip decorated with CNTs and a mercury droplet, demonstrating the role of Hg-nanocontacts. When performed ex situ, this experiment would have wrongly attributed the measured conductance to the ballistic nature of the transport properties of CNTs.
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TEM study of nanostructured BZO templates in (001)-LAO and (001)-STO substrates for the growth of superconducting YBCO films P. Abellan1, M. Gibert1, F. Sandiumenge1, M.J. Casanove2, T. Puig1 and X. Obradors1 1. Institut de Ciència de Materials de Barcelona, CSIC, 08193 Bellaterra, Spain 2. CEMES, 29 rue Jeanne Marvig, BP 94347, 31055 Toulouse Cedex 4, France [email protected] Keywords: YBCO, BZO, CSD
A methodology for the preparation of artificially nanostructured oxide templates by means of chemical solution deposition has already been reported elsewhere [1]. A fine tuning of nanodots’ size, morphology and density has been achieved through the control and modification of precursor solution concentration and growth conditions (temperature, time and atmosphere). Here we apply this methodology to BaZrO3 (BZO). The interest of this study is twofold: first, from the fundamental point of view, it offers a new scenario for the investigation of self-assembling phenomena of oxide nanodots grown from chemical solutions and secondly, because such nanostructured surfaces can be used as templates for the growth of optimized YBCO films. In this work, the microstructure of self-assembled BZO nanoislands generated following a very general bottom-up approach, based on the deposition of chemical solutions, has been investigated using cross-section transmission electron microscopy. TFA-derived YBCO has been deposited on the nanostructured template and the effects of the nanodots on the YBCO lattice has been studied. Two samples have been investigated: one deposited from 0.03M solutions and annealed at 900ºC for 4h in O2, and the other one deposited from a 0.003M solution. On one hand, low magnification TEM images of the 0.03M sample reveal a pronounced bimodal size distribution of particles. Larger particles have dimensions within the range 300 to 400nm and outcrop from within a nanocrystalline film homogeneously covering the substrate. A higher magnification image of the nanocrystalline fraction of the film shows that it is composed by nanocrystals of a tetragonal Ba-Zr oxide phase. Selected area electron diffraction patterns featured very weak diffraction rings consistent with a completely random orientation of this nanocrystalline fraction of the film. On the other hand, for the most diluted precursor solution (0.003M), there is an evolution of the system from the polycrystalline, randomly oriented microstructure to well-separated cube on cube epitaxial cubic BZO nanodots. In this case, BZO grown on LAO and STO substrates present a narrow size distribution of 10 - 20 nm in weight and 5 - 10 nm in height, which has been determined by AFM and confirmed by TEM. Furthermore, AFM analysis reveals BZO nanodots are basically localized close to step edges. HRTEM analysis reveals that BZO forms a self assemblage of well separated cube-on-cube epitaxial islands and are commonly facetted {100} and {110} planes, Figure 1.
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Figure 1. (a) High resolution TEM image of a BZO nanodot grown on LAO single crystal from a solution having a concentration of 0.003M; the nanodots are mainly faceted {100} and {110} planes. (b) AFM image showing a narrow size distribution of the self-assembled nanodots. Observation of strain fields associated with interfacial defects indicates that the BZO nanodots are fully relaxed.
Figure 2. High resolution TEM image of a BZO nanodot at the interface between the YBCO film and the LAO substrate. A high resolution TEM image of an interfacial BZO nanodot, corresponding to a template where YBCO has been deposited, is shown in Figure 2. The image clearly reveals strong bending of the (001) planes at the interface with the nanodot, demonstrating that the presence of the nanodots may disturb the YBCO lattice, while no threading dislocations originating from the nanodots have been identified yet. Epitaxial YBCO quality is preserved in the presence of interfacial BZO dots but, however, in standard growth conditions, the presence of interfacial nanodots enhances the nucleation of a-oriented domains. 1. 2.
M. Gibert, T. Puig and X. Obradors, Surf. Sci. 601 (2007), 2680 This work has been supported by EU (HIPERCHEM project), Generalitat de Catalunya (2005-SGR-0029 and CeRMAE) and by Spanish MEC (MAT2005-02047, FPU program)
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Hydrothermal synthesis and characterisation of single crystal α-Fe2O3 nanorods T. Almeida, Y.Q. Zhu and P.D. Brown School of Mechanical, Materials and Manufacturing Engineering, The University of Nottingham, University Park, Nottingham, NG7 2RD, United Kingdom [email protected] Keywords: hydrothermal synthesis, α-Fe2O3, β-FeOOH, nanorod
One-dimensional (1D) nanostructures have attracted extensive attention due to their novel magnetic properties which are greatly dependent on nanorod size and shape [1]. If the particle size decreases toward a critical value it can consist of a single magnetic domain. Further, a 1D magnetic nanostructure can exhibit much higher coercivity than its isotropic counterpart because of the effect of shape anisotropy. Ferromagnetic hematite (α-Fe2O3) is of particular interest because of its low cost, high resistance to corrosion and environmentally friendly properties. Hydrothermal synthesis (HS) is a comparatively simple process for the preparation of 1D iron oxide nanostructures. However, the mechanism of nanorod formation is not clearly understood, thereby preventing the effective control of α-Fe2O3 nanorod size, shape and growth. α-Fe2O3 nanorods offer many opportunities with applications in fields such as drug delivery, gas sensors and magnetic data storage. A systematic experimental programme has been undertaken to investigate the effect hydrothermal processing conditions on the controlled production of 1D α-Fe2O3 nanostructures. At low hydrothermal temperatures, small β-FeOOH nanorods nucleate from an initial aqueous FeCl3 solution, as shown in Figure 1a. With increasing hydrothermal temperature, an enhancement in the atomic vibration of β-FeOOH molecules results in the breaking of hydrogen bonds, and through the loss of OH- ions [2], the nanorods agglomerate and transform into larger, more thermally stable α-Fe2O3 nanoparticles. The aspect ratio of the nanoparticles may be controlled by the addition of a phosphate surfactant [3]. The PO43- anions have a marked effect on the manner in which small β-FeOOH nanorods align before crystallisation (Figure 1b). The surfactant, however, restricts the dehydration of β-FeOOH and hence suppresses phase transformation into α-Fe2O3. Thus, at higher PO43- concentrations, higher reaction temperatures are required to produce single crystal lenticular α-Fe2O3 nanorods with high aspect ratio (Figure 2). A thorough exploration of the HS growth conditions has been performed using the combined complementary characterisation techniques of transmission electron microscopy (TEM), selected area electron diffraction (SAED), Fourier transform infra-red (FTIR) spectrometry, X-ray diffractometry (XRD) and Xray photoelectron spectrometry (XPS). Through clear understanding of the growth mechanism, it becomes possible to control the size and shape of these 1D α-Fe2O3 nanorods. The importance is made clear
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when α-Fe2O3 becomes super-paramagnetic below certain dimensional values, making it ideal for contrast media applications, e.g. in magnetic resonance imaging [4].
Figure 1. TEM micrographs showing (a) small β-FeOOH nanorods produced through HS at 120˚C and (b) the alignment of β-FeOOH nanorods through the addition of a phosphate surfactant.
Figure 2. A single crystal lenticular α-Fe2O3 nanorod produced through HS at 200˚C. 1. 2. 3. 4.
G.Bate, J. Magn, Magn. Mater. Vol 100 (1991) 413 Y. Sui, D. Xu and W.Su, Materials Research Bulletin, Vol 30, No 12 (1995) 1553-1560 T. Sugimoto, Y. Wang, H. Itoh and A. Muramatsu, Colloids Surf. A Vol 134 (1998) 265 Lawaczeck.R, Menzel.M, Pietsch.H, Applied Organometallic Chemistry, Vol 18 (2004) 506
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GaAs NWs and Related Quantum Heterostructures Grown by Ga-Assisted Molecular Beam Epitaxy: Structural and Analytical Characterization J. Arbiol1,2, S. Estradé2, F. Peiró2, J.R. Morante2, C. Colombo3, D. Spirkoska3, G. Abstreiter3 and A. Fontcuberta i Morral3 1. TEM-MAT, Serveis Cientificotècnics, Universitat de Barcelona, E-08080 Barcelona, CAT, Spain 2. EME/CeRMAE/IN2UB, Departament d’Electrònica, Universitat de Barcelona, C/ Martí i Franquès 1, E-08080 Barcelona, CAT, Spain 3. Walter Schottky Institut, Technische Universität München, D-85748 Garching, Germany [email protected] Keywords: GaAs NWs, MBE, heterostructure, TEM, HAADF, EELS.
Molecular Beam Epitaxy has been for the last two decades the most established technique for the realization of atomically precise nanostructures such as quantum wells, wires and dots. To date the epitaxial precision has been limited to the plane of growth. On the other hand, semiconductor nanowires constitute promising building blocks for next generations of electronic and optoelectronic devices. The use of Molecular Beam Epitaxy (MBE) for the synthesis of ultra-pure 1-D nanostructures has been recently demonstrated. In particular, GaAs nanowires along the [111] direction have been grown avoiding the use of Au as nucleation seeds [1]. The use of the MBE allows us to control the growth conditions during the synthesis and thus being able to tune from 1D to 2D deposition. This latter option led us obtain high quality coaxial GaAs/AlGaAs quantum heterostructures with a perfect epitaxy between the different material layers [2]. In the present work, we will show the structural, morphological and elemental characterization of ultra-pure GaAs NWs and GaAs/AlGaAs coaxial quantum heterostructures grown by molecular beam epitaxy avoiding catalytic seeds. In order to perform the high detail characterization at atomic scale we will use high resolution transmission electron microscopy (HRTEM) and scanning transmission electron microscopy (STEM) in bright field and Z-contrast modes. Electron energy loss spectroscopy will be used to obtain the compositional profiles and maps, which will allow us to distinguish the sharpness in composition of the quantum heterostructures. Special emphasis will be focused on the epitaxial relationships between the different layers of the quantum heterostructures and also on the orientation of the nanowires versus the substrate [2]. The possible crystalline defects will be analyzed (such as lamellar twinning [3,4,5,6]) and the effects of the substrate orientation (Si [001] or [111]) on the nanowire growth and homogeneity of the designed quantum wells commented.
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A.F.I. Morral, C. Colombo, G. Abstreiter, J. Arbiol, J. R. Morante, Applied Physics Letters, 92 (2008), 063112. A.F.I. Morral, D. Spirkoska, J. Arbiol, J. R. Morante and G. Abstreiter, Small, (2008) in press. J. Arbiol, B. Kalache, P. Roca i Cabarrocas, J. R. Morante, A.F.I. Morral, Nanotechnology, 18 (2007), 305606. J. Arbiol, A.F.I. Morral, S. Estradé, F. Peiró, B. Kalache, P. Roca i Cabarrocas, J. R. Morante, Nanotechnology, (2008) in press. F.M. Davidson, D.C. Lee, D.D. Fanfair, B.A. Korgel, J. Phys. Chem. C 111 (2007), 29292935. L.S. Karlsson, K.A. Dick, J.B. Wagner, J.O. Malm, K. Deppert, L. Samuelson, L.R. Wallenberg, Nanotechnology, 18 (2007), 485717.
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Figure 1. BF-STEM and HAADF Z-contrast STEM analysis of a coaxial GaAs/AlAs quantum well heterostructure, viewed along the NW growth axis.
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A method for in-situ electrical measurements of thin film heterostructures using TEM and SEM J. Börjesson1, A. Kalabukhov2, K. Svensson1 and E. Olsson1 1. Microscopy and Microanalysis, Chalmers University of Technology, SE-41296 Göteborg, Sweden 2. Microtechnology and Nanosience, Chalmers University of Technology, SE-41296 Göteborg, Sweden [email protected] Keywords: In-situ TEM, In-situ SEM, STM
A method for in-situ measurements of electrical properties of thin film heterostructures using TEM and SEM has been developed. This method allows measurements of the conductance in a thin film heterostructure in a direction along the film planes in thin TEM foils. The advantage is that the properties can be directly correlated to the local atomic structure. The specimens are cross section TEM samples prepared with standard grinding, polishing and ion beam milling techniques. The equipment used for the measurements is a scanning tunnelling microscope (STM) that can be inserted in both a TEM holder and on a SEM stage. The spatial precision of the STM tip motion is on the atomic scale [1]. The STM tip can translated be controlled in x-, y- and z-direction with both a coarse and a fine motion mode. Figure 1 shows a schematic of a typical TEM thin film cross-sectional sample and the instrument setup. A combined focused ion beam (FIB) workstation and a scanning electron microscope (SEM) is used to create a micrometer sized back contact from the specimen to the 3mm TEM aperture. A tailored screening method is used to protect the electron transparent region of the specimen from ion beam radiation in the FIB/SEM instrument. The advantage to first performing the measurements in a SEM is that it offers a larger field of view and additional information about the three dimensional geometry of the sample and the STM tip. It also makes it possible carry out reference measurements on parts of the sample that is not electron transparent. The measurements performed in-situ in a TEM enable higher spatial precision in the positioning STM tip with respect to the sample. This makes it possible to measure the electrical properties of thin film layers and heterostructures (see Figure 2) and also study the influence of interfaces and local defects. The TEM provides additional information about the specimen microstructure. In addition to the information extracted from the SEM. 1. 2.
K. Svensson, Y. Jompol, H. Olin and E. Olsson, Review of Scientific Instruments 74 (2003), p. 4945. Support from Nanofactory Instruments is gratefully acknowledged for technical support.
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Figure 1. Schematic of (a) a standard TEM thin film cross-sectional sample with a Platinum back contact made by FIB/SEM and (b) the experimental setup with the TEM sample and the STM tip. Electrical measurements are possible by controlling the voltage and measuring the current between tip and sample as a function of position on the specimen.
Figure 2. A TEM micrograph of a thin film superlattice of La1-xSrxMnO3 (blue), LaAlO3 (green) and SrTiO3 (yellow) on an NbGaO3 substrate (red). Courtesy to Boschker, University of Twente, for supplying the thin film heterostructure.
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Electron Beam Nanofabrication and Characterization of Iron Compounds K. Furuya1, M. Shimojo1,2, M. Takeguchi1, M. Song1, K. Mitsuishi1 and M. Tanaka1 1. National Institute for Materials Science; 3-13 Sakura; Tsukuba, 305-0003 JAPAN 2. Advanced Science Research Laboratory; Saitama Institute of Technology; 1690 Fusaji; Fukaya, Saitama, 369-0293, JAPAN [email protected] Keywords: electron beam nanofabrication, iron antenna, iron oxide nanostructures
The production of nanometer-sized oxide structures is of significant interest due to the potential application to nanoelectronics, nanomagnetics and optoelectronics. It is important to place nano-dots or nano-wires at desired positions for the production of nanodevices. Electron beam-induced deposition (EBID) is a promising technique to produce such dots and wires [1]. The resolution of EBID is now reaching down to subnanometers [2]. Recently, Shimojo et al. reported that carbon-free crystalline iron oxide was deposited at room temperature by EBID [3]. In this paper, effects of water and oxygen addition to iron pentacarbonyl on nanostructure are studied. EBID was carried out in a 30kV FE-SEM (JSM-7800UHV) with a custom made gas introduction system. A schematic illustration of the deposition system is shown in Figure 1 [3]. The electron beam current was 0.8 nA with a beam diameter of 4 nm. The base pressure of the chamber was 2x10-6 Pa. Two cylinders, either 1) iron pentacarbonyl (Fe(CO)5) and water (H2O) or 2) Fe(CO)5 and oxygen (O2) were connected. The electron beam position was controlled by an external deflection with a PC. A high resolution TEM image and a diffraction pattern of nanorods formed using a mixture of Fe(CO)5 and H2O vapors are shown in Figs. 2(a) and 2(b). The partial pressure ratio [H2O]/[Fe(CO)5] was 1.0. Crystalline Fe3O4 iron oxide was seen to be formed, by analyzing the diffraction pattern. No amorphous carbon is observed in the nanorods. It is suggested that oxygen radicals formed from water molecules remove carbon and oxidize iron. The composition of the deposits as a function of partial pressure ratio, [H2O]/[Fe(CO)5] is shown in Figure 2(c). The carbon content decreased with increasing the partial pressure ratio, and reached almost zero at partial pressure ratio of 1.0 or more. The ratio of Fe and O became constant at partial pressure ratio of more than 1.0. Figure 3 shows a high resolution TEM image of a nanorod formed with Fe(CO)5 mixed with O2 at a partial pressure ratio [O2]/[Fe(CO)5] of 1.0. Fast Fourier transforms of parts of the high resolution image near the surface and inside regions are also shown. Both reveal that small crystals exist near the surface but the inside is amorphous. The mechanism will be discussed in the talk. 1.
M. Shimojo, K. Mitsuishi, A. Tameike and K. Furuya: J. Vac. Sci. Technol. B 22 (2004) 742.
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W. F. van Dorp, B. van Somern, C. W. Hagen and P. Kruit: Nano Lett. 5 (2005) 1303. M. Shimojo, M. Takeguchi and K. Furuya: Nanotechnology, 17 (2006) 3637.
Figure 1. Schematic drawing of a 30 kV FESEM with dual gas introduction system for EBID with a gas of iron carbonyl
Figure 3. High resolution TEM image (a) of a nanorod formed with Fe(CO)5 mixed with O2 at a partial pressure ratio [O2]/[Fe(CO)5] of 1.0. Fast Fourier transforms of parts of the high resolution image near the surface (b) and inside (c) regions.
Figure 2. High resolution TEM image (a) and a diffraction pattern (b) of nanorods formed at a partial pressure ratio of 1.0, and the composition of the deposits as a function of partial pressure ratio (c), using a mixture of Fe(CO)5 and H2O.
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TEM analysis of the chemical gradient in (Zn,Mn)Te/ZnTe nanowires H. Kirmse1, W. Neumann1, S. Kret2, P. Dłużewski2, E. Janik2, W. Zaleszczyk2, A. Presz3, G. Karczewski2, T. Wojtowicz2 1. Humboldt-Universität zu Berlin, Institut für Physik, AG Kristallographie, Newtonstraße 15, D-12489 Berlin, Germany 2. Institute of Physics, PAS, Al. Lotników 32/46, 02-668 Warsaw, Poland 3. Institute of High Pressure Physics, PAS, Sokołowska 29/37, 01-142 Warsaw, Poland [email protected] Keywords: (Zn,Mn)Te, nanowires, analytical TEM
Magnetic semiconductor nanowires (NWs) are subject of intense fundamental research due to their potential applicability in the field of spin-based electronics. In general, NWs are several micrometers long and only a few tens of nanometers in diameter. Our previous work focussed on the understanding of the structural and chemical properties of ZnTe NWs [1-3]. These NWs were grown on GaAs substrates via a vapour-liquid-solid (VLS) growth regime using a Au-based catalyst droplet. The growth of current (Zn,Mn)Te diluted magnetic semiconductor NW structures was based on the same VLS mechanism and performed at 380 °C in a molecular beam epitaxy machine. The orientation of the GaAs substrate was (110). As seen in the SEM images of Figure 1 the length of the (Zn,Mn)Te NWs amounts to about 400 nm after a period of 15 min (Figure 1a). The combination of 15 min (Zn,Mn)Te and 15 min ZnTe growth reveals an overall length of the NWs of about 1.5 µm. Hence, the growth rate of ZnTe part is twice as large than the apparent growth rate of the (Zn, Mn)Te pedestal. In order to analyse the sharpness of the (Zn,Mn)Te/ZnTe interface, analytical TEM was applied. For localizing the interface a series of EDX point spectra was acquired along the NW axis (cf. Figure 2). Comparing the Mn-K signals of positions 2 and 3 (Figure 2b) a decrease is found due to the termination of the Mn molecular beam after 15 min. The interface is located at a position of about 400 nm. The slow exponential decrease of the Mn signal for the positions 3 to 6 hints either to a high content of Mn still present in the catalyst droplet after closing the Mn shutter or to a strong Mn interdiffusion. The decreasing background signal is due to the shape of the NW. Observing the bottommost area around the NWs a compact layer is seen. Zn, Mn and Te are detected in this area. Moreover, Au and Ga, forming the eutectic catalyst present at the tip of each NW, is found in the layer. In agreement to the results for the NWs a decrease of the Mn content is detected from position A to B whereas Zn increases (see Figure 3b). In the visualized area the layer is separated from the GaAs substrate (cf. Fig. 3a), although in general, this layer is in contact with the substrate. Due to the epitaxial relation between NWs and the substrate, the NWs grow along the two 〈111〉 directions sticking out of (110)-oriented GaAs substrate, as visible in all figures.
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E. Janik, J. Sadowski, P. Dluzewski, S. Kret, L.T. Baczewski, A. Petroutchik, E. Lusakowska, J. Wróbel, W. Zaleszczyk, G. Karczewski, and T. Wojtowicz, Appl. Phys. Lett.89 (2006) p. 133114. E. Janik, P. Dluzewski, S. Kret, A. Presz, H. Kirmse, W. Neumann, W. Zaleszczyk, L.T. Baczewski, A. Petroutchik, E. Dynowska, J. Sadowski, W. Caliebe, W. Wojtowicz Nanotechnology 18 (2007) p. 475606. H. Kirmse, W. Neumann, S. Kret, P. Dłużewski, E. Janik, G. Karczewski, T. Wojtowicz, IOP Conf. Series, Proc. 15th Conf. on Microscopy of Semiconducting Materials, April, 2nd to April 5th, 2007, Cambridge, United Kingdom, in print The research was partially supported by the Ministry of Science and Higher Education (Poland) through grants N507 030 31/0735, N515 015 32/0997 and by the Network "New materials and sensors for optoelectronics, information technology, energetic applications and medicine", and by the Foundation for Polish Science through subsidy 12/2007.
Figure 1. SEM images of growth states of (Zn,Mn)Te/ZnTe NWs; left: 15 min growth of (Zn,Mn)Te (after 15 min the main shutter was introduced in front of this substrate), right: 15 min growth of (Zn,Mn)Te and 15 min growth of ZnTe, respectively.
200 nm
Figure 2. Composition analysis of a (Zn,Mn)Te/ZnTe NWR; left: HAADF STEM image, right: EDX intensity of the Mn-K line for the positions given in the image on the left.
Figure 3. Formation of a layer; left: HAADF STEM image with marker of the EDX line scan given on the right.
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Structural and morphological characterization of GaN/AlGaN quantum dots by transmission electron microscopy M. Korytov1, M. Benaissa2, J. Brault1, T. Huault1 and P. Vennéguès1 1. CRHEA-CNRS, Rue Bernard Grégory, Sophia Antipolis, 06560 Valbonne, France 2. CNRST, 52 Bd Omar Ibn Khattab, BP 8027, Agdal, 10102 Rabat, Morocco [email protected] Keywords: HRTEM, HAADF, quantum dots, MBE
In recent years much attention has been devoted to GaN quantum dots (QDs) due to their potential application as light emitting diodes and lasers diodes operating in the ultra-violet range. QDs allow carriers localisation thus diminishing the influence of the dislocations existing in III-nitrides films on light emission efficiency. The majority of GaN QDs are grown on AlN templates. In this study structural and morphological characterization of GaN QDs grown on Al0.5Ga0.5N templates was carried out by transmission electron microscopy. The main advantage of AlGaN alloys over pure AlN substrate is a possibility to be doped, which enables using GaN QDs as a base for light-emitting devices. The samples under investigation were grown by molecular beam epitaxy using ammonia as nitrogen source and they were composed of three buried layers and one surface layer of GaN QDs separated by AlGaN spacers [1]. A series of samples with varying GaN nominal thicknesses was made. All specimens were studied by High-Resolution Transmission Electronic Microscopy (HRTEM) and High Angular Annular Dark Field (HAADF) technique. HRTEM images were taken under (0004) two-beam conditions. For that the sample was tilted about 6º degree out of the [1120]-zone axis along the [1100] direction to minimize number of the incidental beams. The (0002) beam was centred on the optic axis and the images were formed by using only direct (0000) and one reflected (0002) beams to diminish influence of the change of the local sample thickness and the objective lens defocus to HRTEM contrast [2]. Lost of resolution in the (0002) plane does not hamper the analysis since only distances between basal planes need to be measured. HRTEM images were treated by the Geometrical Phase Analysis (GPA) method, which visualize increase of the local lattice parameter (LLP) into GaN QDs against Al0.5Ga0.5N spacers (Figure 1). HAADF technique was used to obtain chemical contrast in the QDs region (Figure 2). Two main results have to be reported. Firstly, above a critical thickness of deposited GaN, we observed a morphological transition between surface QDs which have a pyramidal shape and buried ones which have a truncated pyramid shape as typically reported for GaN/AlN QDs. Secondly, a stress-driven phase separation in the AlGaN spacers with the formation of Al-rich zones above the QD and Ga-rich zone between QDs was found. Such behaviour with alternating Al- and Ga-rich regions may allow a
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spatial correlation of QDs between successive QDs layers for sufficiently thin AlGaN spacers. 1. 2.
T. Hualt, J. Brault, F. Natali, B. Damilano, D. Lefebvre, L. Nguyen, M. Leroux and J. Massies, APL 92, 051911 (2008) D. Gerthsen, E. Hahn, B. Neubauer, V. Potin, A. Rosenauer, and M. Schowalter, Phys. Stat. Sol (c) 0, No. 6, p. 1668-1683 (2003)
Figure 1. A) HRTEM image of single QD and B) LLP colour-coded map obtained by GPA treatment of the previous image.
Figure 2. HAADF image of two QDs. More bright regions correspond to heavier elements.
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Structure and stability of core-shell AuAg nanopartciels Z.Y. Li1, R. Merrifield1, Y. Feng1, J.P. Wilcoxon1, R.E. Palmer1, A.L. Bleloch2, M. Gass2 and K. Sader2 1. Nanoscale Physics Research Laboratory, School of Physics and Astronomy, University of Birmingham, Birmingham B15 2TT, U.K. 2. SuperSTEM Laboratory, STFC Daresbury, Daresbury WA4 4AD, U.K. [email protected] Keywords: STEM, absorption, core-shell nanoparticles
Nanopartciels comprising two different metallic elements offer additional degrees of freedom for altering their physical properties by varying the atomic composition and atomic arrangement. This could potentially enable a wide range of applications for discovering materials with novel physical properties. However, our present understanding of fundamental issues has been hindered by the lack of knowledge of the precise chemical composition and atomic arrangement within a given bimetallic nanoparticle. Here, we report our recent systematic study on structural and optical properties of core/shell structured AgAu nanoparticles of 4-5 nm diameter by a combination of high resolution scanning transmission electron microscopy (STEM) with UV-visible absorption spectroscopy. The nanoparticles were synthesised by first growing a seed nanocrystal (either Ag or Au) followed by solution phase deposition of several generations of overlayers of shell materials (Au or Ag). The particles were passivated with C12SH during the growth process [1]. Figure 1 shows a comparison of absorption spectra taken from toluene solutions containing core/shell AgAu nanoparticles, revealing strong plasmon resonances and discernable differences in plasmon frequency with different growth sequences. However, optical measurements provide average information that may or may not reflect the heterogeneous nature of bimetallic nanoparticles. Therefore, it is extremely important to develop methods having a high spatially-resolved imaging capability for the internal structure to correlate nanostructure with optical properties. Characterising the core-shell structure of AgAu nanoparticles is a challenging task, as both elements have almost identical lattice spacing. We show that, by utilising the large difference in the atomic numbers of Ag (Z=47) and Au (Z=79), high-angle annular dark field (HAADF) imaging using STEM is capable of determining the possible composition modulation within the nanoparticles [2]. The insert in Figure 1 displays typical three-dimensional plots of Z-contrast images (taken by aberration-uncorrected Tecnai F20 STEM) from the corresponding core/shell AgAu particles. Although no atomic lattice is resolved, distinctly different HAADF intensity contrast variation in two core/shell nanoparticle systems is apparent. It is worth of noting that in the present case, the shell thickness is only 1-2 atomic layers, demonstrating the impressive sensitivity of HAADF-STEM based Z-contrast imaging technique. Using the current generation aberration-corrected STEM, one can obtain further details of internal
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structural and chemical information at atomic-scale. Figure 2 displays one of such examples, where atomic arrangement is clear within Ag core/Au shell nanoparticles. To understand the structure of core/shell AgAu nanoparticles, we consider factors controlling the thermodynamics stability of nanoparticles, including surface energy, cohesive energy and strain, etc. All these factors favour for Ag being at the surface or to form alloyed interface with Au. Our study show that ‘inverted’ Ag/Au nanoparticles can be synthesised suggesting that this structure is at least kinetically stable, though the role that the passivated ligands play needs to be taken into account. In summary, we have successfully demonstrated that the HAADF-STEM imaging is a powerful technique to study the internal structures of core-shell nanoparticles and their evolution. This opens up unprecedented opportunities for detailed investigation on structure/property relationship of these individual building blocks of potential novel functional materials 1. 2.
Z.Y. Li, J.P. Wilcoxon, F. Yin, Y. Chen, R.E. Palmer, R.L. Johnston, Faraday Discussion, 138 (2008) 363-373. Z.Y. Li, J. Yuan, Y. Chen, J.P. Wilcoxon, R.E. Palmer, Appl. Phys. Lett. 87 (2005) 243103.
Figure 1. Three-dimensional HAADF intensity profiles of single AgAu core-shell nanoparticles with UV-visible absorption spectra obtained from the solution containing corresponding nanoparticles. The atomic composition of Ag and Au is 1:1.
a
b Figure 2. Unprocessed (a) and filtered (b) HAADF images of a Ag(core)/Au(shell) nanoparticle by aberration-corrected STEM at SuperSTEM facility.
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In-situ studies on electrical and mass transport in multi-wall carbon and vanadium oxide nanotubes M. Löffler, T. Gemming, R. Klingeler, B. Büchner IFW Dresden, P.O. Box 270116, 01171 Dresden [email protected] Keywords: nanotubes, in-situ, STM
The determination of electronic and mechanical properties of nanoscale materials often presents a challenge with regards to handling and contacting individual structures rather than measuring bulk-scale average properties. In our setup, an in-situ scanning tunneling microscope (STM) is used inside a transmission electron microscope (TEM) to study the influence of electric fields and biases on these individual nanostructures while at the same time observing mechanical and/or structural changes. Electrical and concomitant mass transport in iron filled multi-wall carbon nanotubes (MWCNT) driven by electromigration has been studied (Figure 1). The necessary switching current densities, reversibility and respective breakdown current densities of the empty MWCNT using biases in the range of 0-10V have been measured (c.f. [1]). Furthermore, deflections of the nanotubes by a Lorentz-force under current flow have been observed (Figure 2). Contact resistances, which are of great influence in 2-point measurement setups, have been reduced by pulsed, current-driven annealing. The same setup has been used to determine resistances of individual vanadium-oxide nanotubes (VOxNT) and respective I-V-curves with good reproducibility. For the first time, ohmic characteristics with electrical resistances down to 40MΩ have been observed on VOxNT (cf. [2]) between two tungsten STM-tips as well as in tube-tube junctions between VOxNT (Figure 3). This work presents a starting point for measuring important electronic and electromechanical properties of filled and/or doped carbon and oxide-based nanoscale materials [3]. 1. 2. 3.
K. Svensson, E. Olin, E. Olsson, Phys. Rev. Lett. 93 (2004), 145901 L. Krusin-Elbaum, D. M. Newns, H. Zeng, V. Derycke, J. Z. Sun, R. Sandstrom, Nature 431 (2004), 672-676 This work is funded by the „Pakt für Forschung“ We want to thank U. Weissker, I. Hellmann for the samples and S. Leger for technical assistance.
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+3.6V
-3.6V
Figure 1. Electromigrative movement of the iron filling inside a double-wall CNT. The switching current density is in the order of 10GA/cm2
Voltage Figure 2. Double image of one and the same MWCNT produced by vibration. The cause for vibration is the change of the Lorentz-force due to an AC-current and the static magnetic field of the objective lens. At high currents (top) the filling clearly extends and is assumed to be in the molten state.
Figure 3. Left: Vanadium-oxide nanotube between two tungsten STM-tips. Right: Junction between two VOxNT of different diameter.
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Electron microscopy of nano-magnesium produced by Inert Gas Condensation for hydrogen storage E. Piscopiello1*, E. Bonetti2, E. Callini2, L. Pasquini2 and M. Vittori Antisari1 1. FIM Department, ENEA, C.R. Casaccia C.P. 2400, 00123 Rome, Italy. * C.R. Brindisi, via Appia km 706, Brindisi, 72100, Italy. 2. Department of Physics, Università di Bologna and CNISM, viale Berti Pichat 6/2, 40127 Bologna, Italy. [email protected] Keywords: Inert Gas Condensation, Magnesium, hydrogen storage.
Several issues have been extensively explored in last years in order to improve the hydrogen sorption properties of magnesium for Hydrogen storage. Among other parameters, the material microstructure and particle size play an important role, even if a detailed description of the elementary processes behind the reaction of Mg with gaseous hydrogen are still far to be completely understood [1]. In order to prepare model system having a well defined structure, the Inert Gas Condensation (IGC) technique is extremely powerful, since it allows the formation of dispersed nanoparticles or clusters even with a core shell structure. TEM investigations have been used to study the morphological and structural properties of ICG Mg particles obtained by varying the He gas pressure (0.02÷2 mbar) in the condensation chamber. To this purpose, a FEI Tecnai G2 30F microscope, equipped with a Schottky field emission source and operating at 300kV has been used. Experimental results show that the gas pressure in the IGC chamber affects both the average particle size and the perfection of the structure, as it can be seen in the images reported in Figure 1 (a, b, c). In all batches, the Mg crystals display often a six-fold symmetry which can be correlated with the hexagonal structure of Mg, so that most of the sample is made of (0001) oriented particles as confirmed by the high resolution image reported in Figure 1 (c). Moreover the lateral surface of the particle appears to be decorated with an oxide layer, while the (0001) surface appears to be relatively free from oxygen. The sample produced at the lowest He pressure (0.02mbar, Figure 1(a)) appears to have the smallest average particles size, of the order of a few tens of nm, and a narrower distribution function, while particles synthesized at higher He pressures are larger by about one order of magnitude and with a broad size distribution. High resolution observations showed that in the latter case the particles are single crystal with a perfect structure while in the former a mosaic structure can be observed. The morphological and structural properties of the particles affect the hydrogen sorption behaviour. In fact the particles synthesized at lower pressure show a lower hydrogen desorption temperature (Figure 1(d)) indicating that crystal defects and particle size assists the phase transformation and the hydrogen transport at the free surface.
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In conclusion, the He pressure in the IGC synthesis chamber strongly affects the Mg particle structure with evident effects on the kinetics of reaction with hydrogen. TEM analyses allow an interpretation of the sorption properties on the base of particle microstructure. 1.
A. Zuttel, Materials for Hydrogen Storage, Mater . Today, 6 (2003), pp. 24-33.
(b)
(a)
100nm
(c)
Mg 0,02 mbar
Mg 2 mbar He
200nm
(d)
5nm Figure 1. TEM images of the Mg nanocrystals prepared by IGC at different He pressure, as indicated: 0.02mbar (a) and 2 bar (b). (c) HRTEM image of the border of one particle from the sample synthesized at 2mbar. It is possible to notice the lateral oxide layer and the hexagonal symmetry of the lattice, coherent with the particle shape. Figure 1(d) reports the MgH2 DSC runs, Ar atmosphere, 10 K/min, obtained from particles synthesized at 0.02 and 2mbar of He pressure.
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Two Different Structures of Crystalline Mesoporous Indium Oxide Obtained by Nanocasting Process E. Rossinyol1, E. Pellicer2,3, M. Cabo3, O. Castell1 and M.D. Baro3 1. Microscopy Services, Faculty of Science, Universitat Autonoma de Barcelona, 08193 Bellaterra, Barcelona, Spain. 2. Institut Catala de Nanotecnologia, Campus de la Universitat Autonoma de Barcelona, Edifici Q (ETSE), 08193 Bellaterra, Barcelona, Spain. 3. Departament de Física, Edifici Cc, Universitat Autonoma de Barcelona, ES-08193 Bellaterra, Barcelona, Spain. [email protected] Keywords: mesoporous, indium oxide, nanocasting
Crystalline mesoporous metal oxides with narrow pore size distributions and controllable morphologies are interesting in a wide range of applications such as gas sensing, optics or catalysis. The synthesis pathway known as nanocasting is based in the growth of metal and semi-conductor nanoparticles within the pores of a regular mesoporous material. The use of the regular channels of ordered mesoporous silica as matrices for the controlled growth of nanoparticles and nanorods has been widely reported. In this method, the three-dimensional pore system of the matrix is filled with a suitable metal oxide precursor, which is then converted to the oxide during the calcination process. The structure matrix is finally removed, yielding the metal oxide as its negative replica. This nanocasting method has allowed the synthesis of numerous mesoporous metal oxides, many of which were formerly unavailable by the conventional synthesis method of supramolecular structure direction. Indium is a low melting point metal, which can easily be oxidised into indium oxide (In2O3). Indium oxide is a transparent n-type semiconducting oxide with a band gap of around 3.6 eV, which is known to exhibit luminescence1 and gas sensing2 properties. In addition, indium oxide nanofibers3 and nanosized particles dispersed within pores of non-organised mesoporous silica4 have been reported to give photoluminescent properties. For our aim, two different mesoporous silica structures were chosen: i) the 3D cubic (space group Ia3d) named KIT-6, which is probably the most complex one presenting a double gyroidal mesostructure with channels running along the [100] and [111] directions, defining a 3D pore network;4 ii) the 2D hexagonal structure (space group p6mm) named SBA-15, which has a simpler structure. Both templates have been used for the synthesis of mesoporous indium oxide nanostructures that were compared and have shown evidence of possible structure effects on the electronic properties. Different catalytic additives have been introduced and their effect have been analysed. In2O3 powder was synthesized by impregnation of 0.15 g mesoporous SBA-15 and KIT-6 silica template with 0.39 g In(NO3)3 · xH2O (purchased from Aldrich) in ethanol. Fully details of silica template synthesis were reported elsewhere6. After the
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impregnation step, the resulting mixtures were dried and calcined at 350 °C for 4 h. Then, a second impregnation was carried out by mixing the powder obtained with 0.22 g of indium nitrate precursor. After 30 min of stirring, samples were dried and calcined again at 600 °C for 4 h. Finally, the silica template was removed with NaOH solution. The slurry obtained was centrifuged, copiously washed and dried. Catalytic additives were added by impregnation. Water used had been distilled twice and deionized with a Millipore Milli Q system. In the wide-angle XRD patterns of the In2O3 powders, the main diffraction peaks correspond to cubic Ia3 In2O3 phase (JCPDS card no. 06-0416). Peaks are quite broad, which demonstrates the nanometric-grained nature of the samples. The absence of a wide peak at about 22º (2θ) related to amorphous silica proves the purity of In2O3. TEM images were taken to assess the quality of both templates (Figure 1a and 1b) and the mesoporous oxide replica. Pure In2O3 samples are highly crystalline, with the typical cubic (byxbyite) structure as observed in XRD patterns. Monocrystalline nano-scaled wires of 5-8 nm in diameter and around 200 nm in length can be observed (Figure 1c)
Figure 1 Mesoporous silica used as a template: KIT-6 a) and SBA-15 b). Detailed of an indium oxide replica of SBA-15 structure. 1. 2. 3. 4. 5. 6.
M.-S. Lee, W. C. Choi, E. K. Kim, C. K. Kim and S.-K. Min, Thin Solid Films, 279 (1996) 1 Prim, E. Pellicer, E. Rossinyol, F. Peiró, A. Cornet and J. R. Morante. Advanced Functional Materials, 17 (2007) 2957. Liang, G. Meng, Y. Lei, F. Phillipp and L. Zhang, Advanced Materials, 13 (2001) 1330. Murali, A. Barve, V. J. Leppert, S. H. Risbud, I. M. Kennedy and H. W. H. Lee, Nano Lett, 1 (2001) 287. Tian, X. Liu, L. Solovyov, Z. Liu, H. Yang, Z. Zhang, S. Xie, F. Zhang, B. Tu, C. Yu, O. Terasaki, D. Zhao. J. Am. Chem. Soc. 126 (2004) 126. Zhao, J. Feng, Q. Huo, N. Melosh, G. H. Fredrickson, B. F. Chmelka, G. D. Stucky, Science 279 (1998) 548
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Measuring electrical properties of carbon nanotubes using liquid metal immersion, an in situ scanning electron microscopy study H. Strand1, K. Svensson2 and E. Olsson3 1. Department of Physics, Göteborg University, SE-412 96 Göteborg, Sweden 2. Department of Physics and Electrical Engineering, Karlstad University, SE-651 88 Karlstad, Sweden 3. Department of Applied Physics, Chalmers University of Technology, SE-412 96 Göteborg, Sweden [email protected] Keywords: Liquid metal, carbon nanotube, SEM
The electrical properties of nano-materials are of great interest for both science and industry. In particular carbon nanotubes [1] are one of the most studied materials. It is known that the electrical properties of single-walled carbon nanotubes are determined by their chirality [2] and can be either metallic or semiconducting. Experiments have shown the electron transport to be ballistic at room temperature [3]. In multi-walled carbon nanotubes (MWCNTs) the electron transport is more complicated due to the various layers of different chirality, and the dominating transport regime at room temperature is still under some debate. Early low temperature magnetoresistance measurements found evidence of weak localization [4], a correlation effect arising in diffusive systems. These works predicted the transport to be diffusive also at room temperature, as was later confirmed by scanning probe microscopy [3]. All these measurements were performed on substrate deposited MWCNTs. Another very different and exciting experimental approach for room temperature measurements has been introduced by S. Frank et. al. [5]. They mounted MWCNTs on a piezo-controlled tip and used a liquid metal as counter electrode. Immersing a tube by moving the tip then enabled control of the conducting tube length. Applying a voltage (≈100mV) and measuring the current when moving the tip then gave the resistance R as a function of tube length. In the final report [6] it was concluded that the as-produced arc discharge MWCNTs used did display ballistic conduction and electronic mean free paths of the order 10μm. Spurred by these astonishing results we have performed a study of the liquid metal immersion of individual MWCNTs both in air and in situ in a Zeiss Ultra 55 FESEM using an in house developed scanning tunneling microscope (STM) probe. The probe is based on the same principle as the transmission electron microscope (TEM) holder described in [8], now commercially available from [7]. Our MWCNTs was grown by pyrolysis of ferrocene [9] and were known to be diffusive conductors from previous works [10]. We used the model presented in [6] to interpret the conductance traces from our immersion experiments in air. A typical long trace is shown in “Figure 1a-b”. The
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obtained resistance per length, ρL, was surprisingly low with, ρL≤120Ω/μm, which is of the same order as previous ballistic claims (ρL≤48Ω/μm [6]). For proper interpretation of our data we repeated the immersion experiments in situ in the SEM. The topological sensitivity of secondary electron imaging enabled us to directly image the liquid metal surface response as the MWCTN was pushed in to it. A SEM-micrograph of the surface is shown in “Figure 1c”, where the bright lower surface is the liquid metal (Hg), showing the formation of an inwards meniscus with a large radius of curvature compared to the length of the MWCNT protruding from the upper, movable, tip. The non-wetting meniscus formation makes the model-calculations invalid, since the model assumes that the tube immersion depth is directly given by the piezo-motion of the tip. We have also found evidence of the formation of a thin layer on the liquid metal surface, a known problem for mercury called “skinning” [11]. The immersion model of [6] seems to have gained acceptance and are used in more recent works such as [12]. Thus we want to stress the importance of monitoring the liquid metal surface response to enable correct modeling of the liquid metal immersion experiment. For this purpose the simultaneous immersion and in situ imaging using SEM has proven to be the ideal tool.
c) 1μm
Figure 1. Electrical measurements of a) the conductance G (G0=2e2/h≈1/12.9kΩ) and b) resistance R as function of piezo movement x, linear fit in b) gives ρL, c) meniscus formation at the liquid metal surface. 1 2 3 4
Iijima, Nature 354 (1991), p. 56-58. M. S. Dresselhaus, G. Dresselhaus and R. Saito, Carbon 33 no. 7 (1995), p. 883-891. A. Bachtold et. al., Phys. Rev. Lett. 84 no. 26 (2000), p. 6082-6085. L. Langer et. al., Phys. Rev. Lett. 76 no. 3 (1996), p. 479-482; A. Bachtold et. al., Nature 379 (1999), p. 673–675; C. Schönenberger et. al., Appl. Phys. A 69 (1999), p. 283-295. 5 S. Frank et. al., Science 280 (1998), p. 1744-1746. 6 P. Poncharal et. al., J. Phys. Chem. B 106 (2002), p. 12104-12118. 7 Nanofactory Instruments AB, www.nanofactory.com. 8 K. Svensson et. al., Rev. of Sci. Instr. 74 no. 11 (2003), p. 4945-4947. 9 Produced by H. Pettersson, M. Terrones and N. Grobert, Sussex Nanoscience and Nanotechnology Centre, University of Sussex, England. 10 K. Svensson, H. Olin and E. Olsson, Phys. Rev. Lett. 93 no. 14 (2004). 11 M. C. Wilkinson, Chem. Rev. 72 no. 6 (1972), p. 575-625. 12 H. Kajiura et. al., Chem. Pys. Lett. 398 (2004), p. 476–479; Carbon 43 (2005), p.1317–1339; Appl. Phys. Lett. 86 (2005) art. 122106
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TEM characterization of biogenic metal nanoparticles E.I. Suvorova1,3, P.A. Buffat1, H. Veeramani1, J. Sharp1, E. Schofield2, J. Bargar2, and R. Bernier-Latmani1 1. Environmental Microbiology Laboratory, Ecole Polytechnique Fédérale de Lausanne, Station 6, CH-1015 Lausanne, Switzerland 2. Stanford Synchrotron Radiation Laboratory, Menlo Park, CA 3. Institute of Crystallography RAS, Leninsky pr., 59, Moscow 119333, Russia elena.suvorova @epfl.ch Keywords: Cr and U oxides nanoparticles, microbial reduction, HRTEM and X-ray EDS analysis
Biogenic nanoparticles formed by microbial reduction of uranyl and chromate were investigated in order to understand the mechanism and to design strategies for the bioremediation of Cr (VI) and U (VI) contaminated groundwater. Analysis of composition, structure, morphology and sizes of particles were carried out in a FEI CM300UT FEG electron microscope in TEM/HRTEM/Electron diffraction/STEM/EDS mode. The images were recorded on a Gatan 797 slow scan CCD camera and processed with the Gatan Digital Micrograph 3.11.1 software, INCA (Oxford) and JEMS package [1]. Low intensity illumination conditions were used to record the HRTEM images. Element-specific short-range structural and chemical information about the U coordination environment was derived from the Extended Xray Adsorption Fine Structure (EXAFS) spectra collected at room temperature or at 77K at the Stanford Synchrotron Radiation Laboratory and processed with the software package SIXPACK. The microbial reduction of uranyl and chromate, the toxic and highly soluble forms of the metals, to sparingly soluble Cr (III) as Cr (III) hydroxide/oxide/phosphate and U(IV) as UO2 was observed. Uraninite, UO2, a face-centered cubic mineral, occurred as nanoparticles, (Figure 1). Their size did not exceed 4.0±0.2 nm in diameter. The addition of Mn or Mg cations during the reduction process resulted in a reduction in uraninite particle size: their diameter did not exceed 3.0 ± 0.2 nm. Cr(VI) reduction was studied in the presence and absence of Fe(III) and yielded a non-crystalline precipitate. STEM and EDS methods were used to characterize this material. Previous work has shown that a mixed FexCry(OH)3 phase was formed when both Cr(VI) and Fe(III) were present [2]. The present STEM/EDS data support two observations: (i) a Cr(III)-only phase was associated with the bacterial surface (Figure 2a) and (ii) elemental maps of particle aggregates away from the bacterial surface suggested that several Cr and Fe phases, for instance CrPO4, Cr2O3, Cr(OH)3, FePO4 and Fe2O3 could form and coexist (Figure 2b). 1. 2.
P. Stadelmann, JEMS. http://cime.epfl.ch, 2008. C.M. Hansel, B.W. Wielinga, S. Fendorf. Geochimica et Cosmochimica Acta, Vol. 67 (2003), pp. 401–412.
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Figure 1. TEM (a), SAED (b), HRTEM (c), and a diffractogram (d) taken from the UO2 nanoparticles. Figure 2. STEM images and element maps of Cr (III) precipitates on the bacterial surface (a) and an agglomerate with mixture of Fe (III) and Cr (III) particles (b).
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On the structure of VxOy supported on multiwalled carbon nanotubes D. Wang1, J.-P. Tessonnier1, M. Willinger1,2, C. Hess, D.S. Su1, R. Schlögl1 1. Abteilung Anorganishe Chemie, Fritz-Haber-Institut der Max-Planck-Gesellschaft, Faradayweg 4-6, D-14195 Berlin, Germany 2. Department of Chemistry, CICECO, University of Aveiro 3810-193 Aveiro, Portugal [email protected] Keywords: Vanadium oxide, catalysis, carbon nanotubes, TEM
Supported highly dispersed Vanadium oxides are of great interests for the high activity in selective oxidation of small alkanes (propane, butane) [1]. CNTs are widely used in heterogeneous catalysis as support as it is possible to manipulate them by specific chemical and thermal treatments aimed to tune the surface chemistry [2]. By incipient wetness impregnation with vanadium oxide triisopropoxide in isopropanol, highly dispersed vanadium oxide on commercial carbon nanotubes (Baytubes) with loadings of 3, 10 and 20 wt. % V were synthesized. Catalytic reactions were performed on small amounts of catalyst directly mounted onto a gold TEM grid in a microreactor, allowing direct correlation between structure and catalytic behaviors. The catalysts with lower vanadium loading show a high initial activity at 400°C but deactivate quickly. The VxOy/MWCNTs samples with different loadings were investigated by (HR)TEM, EDX ,EELS, XPS and Raman spectroscopy. The HRTEM image from 3 wt.% VxOy/CNTs (Figure 1a) looks very similar to the pure CNTs, while the EDX spectrum from an area of about 1 μm2 shows confirm the V quantity (2.8 wt. %) in agreement with the nominal value. The observation indicates two-dimensionally dispersed VxOy species both inside and outside of the CNTs. With increase of V loading to 10 wt. %, the morphology of CNTs doesn’t change noticeably and the EDX spectrum from large area confirms the V loading (12.9 wt.%). However, aggregated islands are more frequently observed, as shown in Figure 1b with corresponding local EDX spectrum in the inset. These islands are mostly flattened and extend along CNTs surfaces with the thickness of 1-2 nm. Such aggregation becomes more prominent for 20 wt. % VxOy/CNTs. V oxides form continuous aggregates over large area on the surface of CNTs or fill the inner space within the CNTs. After test in oxidation of propane to propylene at 400°C, the CNTs support was destroyed for all the catalysts and the TEM image shows that a considerable amount of V oxides form crystals. Representatively, Figures 2a, 2b and 2c show the overview image, selected area diffraction pattern and the HRTEM image of the 10 wt. % VxOy/CNTs after the reaction, respectively. The HRTEM image shows the lattice spacing other than V2O5 while the diffraction pattern is in good agreement with the simulated V2O5 diffractions (inset in Figure 5b). Such result implies that the main structure of the V oxides after the reaction is V2O5 and some suboxides phases may coexist on the surface. S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 317–318, DOI: 10.1007/978-3-540-85226-1_159, © Springer-Verlag Berlin Heidelberg 2008
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The oxidation states of V oxides for as prepared catalysts and those after the reaction were also studied by EELS. From the ELNES of V L2,3 edge and O Kedge, an oxidation state lower than 5+ can be deduced for the 3 % V catalyst. The as prepared catalysts with higher loading have varying V oxidation states in different areas. After test in oxidation of propane, the aggregated V oxides show exclusively an oxidation state close to 5+, in agreement with the electron diffraction result. In summary, highly dispersed VxOy species are highly active in selective oxidation of small alkanes. Deactivation is attributed to the formation of bulky vanadium oxides. How to stabilize the VxOy species is essential for optimization of the catalysts. 1 2 3
C. Hess, U. Wild and R. Schloegl, Microporous Mesoporous Mater. 95 (2006), p. 339. W. Xia, Y. Wang, R. Bergstraesser, S. Kundu and M. Muhler, Appl. Sur. Sci. 254 (2007), p. 247. The authors acknowledge financial support by SFB 546 “Transition Metal Oxide Aggregates.”
Figure 1 HRTEM images of VxOy/MWCNTs a) with a loading of 3 wt% V and b) 10 wt% V (right). The formation of VxOy islands is confirmed by EDX.
Figure 2 a) Overview image, b) selected area diffraction pattern together with the simulated pattern and c) HRTEM image of the 10 wt% V/MWCNT catalyst after reaction, respectively.
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Hexahedral nano-cementites catalysing the growth of carbon nanohelices J.H. Xia1,3, X. Jiang1,3, C.L. Jia2 and C. Dong3 1. Institute of Materials Engineering, University of Siegen, Paul-Bonatz-Str. 9-11, 57076 Siegen, Germany 2. Institute of Solid State Research and Ernst Ruska Centre for Microscopy and Spectroscopy with Electrons, Research Centre Jülich, D-52425 Jülich, Germany State 3. Key Laboratory of Materials Modification & Department of Materials Engineering, Dalian University of Technology, Dalian 116024, P. R. China [email protected] Keywords: Carbon nanohelices, hexahedral nano-cementites, growth mechanism
Helical carbon structures (fibers, tubes, etc) have attracted considerable attention of research owing to their attractive morphology and properties [1,2]. Carbon nanohelix (CNH) is one type in the helical family, which has been grown on iron needles using microwave plasma assisted chemical vapor deposition (MPCVD) [3]. To explore the formation mechanism of CNH, systematic investigations on the catalyst particle are carried out by means of scanning electron microscopy (SEM) and transmission electron microscopy (TEM). The catalyst is identified to be single crystalline cementite (Fe3C) particle at the tip of each CNH. The Fe3C particles show a general morphology of hexahedron with six (different) crystallographic planes as the surface planes, which is determined by shape fitting with the series of TEM images (Figure 1.). The different catalytic effect of different crystallographic surface planes produces an anisotropic growth on the front surface of the carbon nanostructure, which results in a rotation of the cementite particles. The rotating particles catalyse the growth of the carbon nanostructure in a helix way. The schematic growth model is shown in Figure 2.
1. 2. 3.
S. Amelinckx, X. B. Zhang, D. Bernaerts, X. F. Zhang, V. Ivanov and J. B. Nagy, Science 265, 635 (1994). S. M. Yang, X. Q. Chen and S. Motojima, Appl. Phys. Lett. 81, 3567 (2002). G. Y. Zhang, X. Jiang, and E. G. Wang, Appl. Phys. Lett. 84, 2646 (2004).
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Figure 1. (Up left) Series of TEM images and corresponding diffraction patterns of Fe3C catalyst particles recorded along different crystallographic directions. In (a)-(d) the images and diffraction patterns are obtained by tilting a particle around the [110] axis (arrows in diffraction patterns). In (e) and (f) the images and diffraction patterns are recorded by tilting around the [001] axis. Details are described in main text. A hexahedron model of the particle is shown in (g) and its projections are shown accompanying the image series. Figure 2. (Up right) The schematic growth model of the nanohelices.
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Investigation of organic/inorganic interfaces using nano-analytical transmission electron microscopy V. Jantou1, M.A. Horton2 and D.W. McComb1 1. Dept. of Materials, Imperial College London and London Centre for Nanotechnology 2. Dept. of Medicine, University College London and London Centre for Nanotechnology. [email protected] Keywords: EELS, FIB milling, apatite, collagen
The impact of long-term health problems due to skeletal disorders, such as osteoporosis and arthritis, is increasing with the ageing of modern populations. Bone pathologies have a critical effect on the mechanical properties of bone, which are intimately linked to its highly complex hierarchical structure. However, despite decades of investigation, bone nanostructure (essentially made of collagen fibrils and apatite crystals) is not well understood. The exact structure and chemistry of bone crystals and their location with respect to collagen is still debated. The currently accepted 2D model [1] states that the collagen molecules are arranged in a quarter-stagger fashion, creating gap and overlap regions; the crystals are thought to be located inside the gap zone, giving rise to a periodic contrast in electron micrographs of bone collagen. In order to gain a better understanding of the relationship between the apatite and collagen, it is crucial to study the composition and bonding along collagen fibrils. In the present work, the organic/inorganic interface in mineralised ivory dentin, which acts as a good model for bone, was studied using analytical scanning transmission electron microscopy (STEM), with samples prepared using focused ion beam (FIB) milling as an alternative technique to ultramicrotomy. A monochromated FEG(S)TEM (FEI Titan 80-300) instrument was used, providing both high-resolution imaging and spectroscopy capabilities using electron energy-loss spectroscopy (EELS) and energy-filtering TEM (EFTEM). The Titan offers a spatial resolution <0.14 nm and an energy resolution <0.2 eV. A clear advantage of FIB milling over ultramicrotomy is that dehydration, embedding and section flotation can be obviated, thus reducing both physical and chemical damage to the specimen prior to examination. FIB milling has other significant advantages over ultramicrotomy, such as site-specificity and the ability to image whilst milling, which provides greater control over the final section thickness. The characteristic periodic contrast of collagen fibrils was visible in FIB sections without the need for any chemical staining, contrary to ultramicrotomed sections. Consequently, the origin of the collagen periodic contrast can now be properly investigated by performing compositional analysis on the nanometer scale, without the chemical complications arising from the sample preparation procedure. EELS and EFTEM mapping of P, O and Ca, the principal constituents of apatite crystals, and C, the major component of the organic phase, are currently being carried
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out, and data is acquired along and across collagen fibrils. By studying not only the position and shape of a peak, but also the fine structure after each edge onset, termed electron energy-loss near-edge structure (ELNES), the atomic environment of a particular species can be determined [3], allowing differently bonded Ca or C atoms to be distinguished. It is the first time that some of the fundamental questions of mineralised tissues can be approached using nano-analytical TEM, which will ultimately provide crucial information regarding the mechanisms of bio-mineralization and how it is altered in diseases. 1. 2. 3.
Weiner S. and Traub W., FASEB Journal 6 (1992), p. 879. Egerton, R.F., NY-London Plenum Press (1996). The authors gratefully acknowledge the Department of Materials, Imperial College London, and University College London (Dr. Mortimer and Mrs. Theresa Sackler Trust).
Figure 1. BF-TEM images of FIB milled and ultramicrotomed sections along (long.) and across (trans.) collagen fibrils. The dotted encircled region defines the crosssectional profile of a collagen fibril.
Figure 2. STEM-EELS line profile across a collagen fibril on a FIB milled section. The spectra were recorded at 100kV, with an energy resolution of 0.6eV, a dispersion of 0.2eV/pixel, and convergence/collection semi-angles of 10.5mrads and 11.8 mrads respectively.
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Cationic ordering and interface effects in superlattices and nanostructured materials P. Boullay, W.C. Sheets, W. Prellier, E.-L. Rautama, A.K. Kundu, V. Caignaert, B. Mercey and B. Raveau Laboratoire CRISMAT, UMR 6508 CNRS/ENSICAEN, 6 Bd Maréchal Juin, 14050 Caen Cedex 4, France [email protected] Keywords: perovskites, HRTEM, heterostructures, domain texture
Thin film methods offer a versatile approach to overcome the natural preference for disorder or low-dimensional ordering in certain materials by controlling the location of cations in tailored oxide multilayers. Artificial heterostructures have hence been utilized to probe how interlayer coupling and cation ordering at the interface affect magnetic or transport properties. Recently, it was shown that LaTiO/SrTiO3 superlattices exhibit metallic conductivity, similar to a bulk solid solution of La1−xSrxTiO3, even though the heterostructure is based on two insulators [1]. In this contribution, a series of highquality (LaVO3)6m(SrVO3)m superlattices (Figure 1) with a nominal composition close to La6/7Sr1/7VO3 has been grown on [001]-oriented SrTiO3 substrates using standard pulsed-laser deposition [2]. When m=1, the bulk-like insulating resistivity of La6/7Sr1/7VO3 is obtained while for m≥2 the superlattices display a metallic behaviour. The comparison of the superlattice periodicity with the known coherence length of charge carriers across the interface of perovskite oxide heterostructures is used to understand this insulator-metal transition. The magnetic and transport properties of manganites and cobaltites with the perovskite structure have been the object of numerous studies. A modification of the cation ordering in such strongly correlated electronic oxides can also affect properties. From this perspective, our recent study of the perovskite La0.5Ba0.5CoO3 has allowed three forms to be isolated. Besides the known disordered La0.5Ba0.5CoO3 and layered 112-type LaBaCo2O6 forms, a third nanoscale-ordered form is obtained [3]. This new form presents a specific nanostructure that consists of 112-type 90° oriented domains fitted into each other at a nanometer scale in order to form a 3D domain texture [4] that induces large strains in the material (Figure 2). Differently from the other forms, this latter is a hard ferromagnet due to the strains which may pin domain walls, preventing the reversal of the spins in a magnetic field illustrating that the atomic-scale lattice distortions are coupled to the magnetic and electronic degrees of freedom. 1. 2. 3. 4.
A. Ohtomo, D.A. Muller, J.L. Grazul and H.Y. Hwang, Nature 419 (2002), 378. W.C. Sheets, B. Mercey and W. Prellier, Appl. Phys. Lett. 91 (2007), 192102. E-L. Rautama, P. Boullay, A.K. Kundu, V. Caignaert, V. Pralong, M. Karppinen and B. Raveau, Chem. Mat. (2008), in press. M.A. Alario-Franco, M.J.R. Henche, M. Vallet, J.M. Gonzalez-Calbet, J.C. Grenier, A. Wattiaux, P. Hagenmuller, J. Solid State Chem. 46 (1983), 23.
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Figure 1. Cross-sectional HRTEM image obtained for a metallic (LaVO3)17(SrVO3)3 superlattice with a periodicity of ~78Å.
Figure 2. Bright-Field images obtained for the nanostructured LaBaCo2O6. A mottled contrast consisting of curved and interpenetrated dark segments is observed and related to the existence of strain fields induced by the 3D texture at a nanometer scale.
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Strain in SrTiO3 layers embedded in a scandate/titanate multilayer system D. Ávila1, M. Boese1, T. Heeg2, J. Schubert2 and M. Luysberg1 1. Institut für Festkörperforschung und Ernst Ruska-Centrum, Forschunszentrum Jülich, 52425 Jülich, Germany 2. Institut für Bio- und Nanosysteme, Forschunszentrum Jülich, 52425 Jülich, Germany [email protected] Keywords: SrTiO3, strain
Deposition of earth alkali titanate thin films on different substrates allows to tune the strain and, as consequence, their dielectric properties. In this sense, the large range of available lattice constants in rare earth scandates make them very useful candidates to serve as substrate materials for epitaxial growth. In this work we have employed aberration corrected high-resolution transmission electron microscopy (HRTEM) to quantify the strain in different scandate/titanate multilayers by using the geometrical phase analysis method [1]. Besides, we applied electron energy loss spectroscopy (EELS) and ab initio calculations to obtain information about the effect of the strain on the electronic structure of these materials. The multilayers of DyScO3/SrTiO3 (DSO/STO) and GdScO3/SrTiO3 were grown by Pulsed Laser Deposition (PLD) using GdcSO3 or DyScO3 substrates [2,3]. The HRTEM study was carried out in an image corrected CM200 microscope. EEL spectra were recorded in a probe corrected Titan microscope. Ab-initio calculations were performed with the Feff8.2 code based on the real space multiple scattering method. Exit-plane wave reconstruction [4] was utilized for the analysis of DyScO3/SrTiO3 multilayers recording a focus series of 20 images with a focal increment of 1.7 nm. The phase image shown in Fig. 1 is analysed with the geometric phase analysis. Linescans corresponding to the red frame in Fig 1 have been performed for the in plane (ip, parallel to the interface), and out of plane (op, growth direction) lattice constants, respectively. Clearly a tetragonal distortion of the STO layers is revealed, which amounts to a maximum 2% of tensile strain, translating to 8 pm smaller lattice parameter of SrTiO3 compared with the in plane lattice parameter of 395 pm= aDSO. EEL spectra have been recorded for the strained STO layers within the multilayer system (red curve in Fig. 2). Compared to a reference spectrum obtained from a STO substrate, a significant smaller splitting of the L3 edge is detected. In cubic SrTiO3 the Ti is surrounded by six O forming a regular octahedron with Ti-O distances 0.195 nm. However, under tensile strain this octahedron is distorted yielding an average distance Ti-O 0.196 nm. This enlargement of the distances produces a weaker crystal field, which results in a smaller splitting. The reduced crystal field splitting is reproduced by ab initio calculations of the Ti L3 edges using the Feff8.2 code (see Fig. 2). Calculations of the L3 edge for different strain situations reveal, that the splitting of the L3 linearly decreases with increasing strain.
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1. 2.
M. J. Hytch, E. Snoeck, R. Kilaas, R. Ultramicroscopy 74 (1998) T. Heeg, M. Wagner, J. Schubert, Ch. Buchal, M. Boese, M. Luysberg, E. Cicerrella, J.L. Freeouf, Microelectronic Engineering 80, (2005) 150 3. M. Boese, T. Heeg, J. Schubert, M. Luysberg, J. Mat. Sci. 41, (2006) 4434 4. A. Thust, W. Coene, M. Op de Beeck, and D. Van Dyck, Ultramicroscopy 64 (1996), 211230 *present address: Carl Zeiss NTS GmbH, Carl-Zeiss Str. 56, 73447 Oberkochen STO
DSO
STO
DSO
STO ip
DSO op
STO STO
DSO
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Figure 1: Phase image obtained from a through focus series of STO/DSO multilayers. The linescans shown on the right hand side correspond to the in plane and out of plane lattice parameters obtained from corresponding geometrical phase images. The in plane lattice parameter shows only small variations, whereas the out of plane parameter varies by 2% across the layer system. Cubic
Intensity (a.u.)
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Ti-O: 1.95
ΔE = 2.22 eV ΔE = 2.01 eV
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ΔE = 2.12 eV ΔE = 2.01 eV
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Figure 2: Experimental and calculated EEL spectra of the Ti L3 edge obtained for cubic SrTiO3 (black curves) and tensilely strained SrTiO3 (red curves). In the strained case, i.e. the in plane lattice parameter resumes the value of DSO, the average Ti-O bond is by 0.01 Å larger, resulting in a smaller crystal field splitting.
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Anisotropic growth of CGO islands on the (001)-LaAlO3 surface A. Benedetti, M. Gibert, F. Sandiumenge, T. Puig and X. Obradors Institut de Ciència de Materials de Barcelona, CSIC, Campus de la UAB, 08193 Bellaterra, Catalonia, Spain [email protected] Keywords: nanohuts, CGO, chemical solution deposition
Chemical solution deposition (CSD) has emerged as a very competitive, low cost and easily scaleable technique for obtaining highly epitaxial multilayered structures, and very recently its high potential for the generation of high quality self assembled templates has been also demonstrated [1]. As a further step, in this contribution we focus on the correlation between shape anisotropy and anisotropic mismatch of Ce1-xGdxOy, x=0.1, (CGO) islands grown on (001)-LaAlO3 (LAO) by CSD, investigated using x-ray diffraction pole-figure analysis, atomic force microscopy (AFM), and high resolution electron microscopy (HREM). Anisotropic mismatch on an isotropic substrate surface in the CGO-LAO system is promoted by the strong sensitivity of island orientation to growth conditions. In particular, we have recently demonstrated that the usual (001)[110]CGO//(001)[100]LAO epitaxial relationship can be replaced by the (110)[1-10]CGO//(001)[100]LAO one [1]. Fig. 1(a) is an AFM image of highly anisotropic nano-huts grown at reduced pO2, exhibiting the (110)[1-10]CGO//(001)[100]LAO epitaxial relationship. This configuration involves an extremely high misfit direction (<100>CGO//<100>LAO, ε~-30%) coexisting with a low misfit direction ([1-10]CGO//[100]LAO, ε ~-1%). Fig. 1(b) shows a HREM cross sectional view of a nano-hut, viewed along the [1-10]CGO direction, i.e., the long direction of nano-huts in Fig. 1(a). The nano-huts are narrower along the higher mismatch strain direction. Despite the huge mismatch, overall the image demonstrates good crystallinity (image simulations indicate that white dots correspond to atom columns of the fluorite-type CGO structure), though the inner region exhibits a change in contrast due to strain localization, very likely associated with the solid-phase growth mechanism of the islands. The interface appears structurally incoherent, which can be understood taking into account the huge mismatch strain. Since the most likely interfacial component of the Burgers vector for CGO is ½<100> [2], full relaxation of such a huge mismatch strain requires a dislocation spacing of 0.9 nm, i.e., a dislocation every other unit cell. This distance agrees well with what we experimentally measured by Fourier analysis (1 - 1.1 nm), indicating that nearly full relaxation takes place. Insights into the strain state within the island where obtained from strain maps obtained by the geometric phase method [3], using an undistorted area of the island as a reference. Images 1(c) and 1(d) display the strain tensor components εxx and εyy, where x and y are parallel and normal to the interface, respectively. With the exception of the central defective region, both strain components, in particular εxx which is more
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sensitive to in-plane strains, reveal a rather homogeneous strain distribution, with values comprised in the range -0.01< ε<0.01 which can hardly be associated with an elastically strained island as a consequence of such a large mismatch. Interestingly, despite the high degree of relaxation the bottom corners of the island (encircled in Fig. 1(b)) show some degree of disorder and are lifted from the substrate surface, suggesting that besides full relaxation within the islands, structural difference with the substrate may represent an additional obstacle to island growth. 1. 2. 3.
M. Gibert, T. Puig, X. Obradors, A. Benedetti, F. Sandiumenge, R. Hühne, Adv. Mater. 19 (2007) pp. 3937–3942 A. Cavallaro, F. Sandiumenge, J. Gazquez, T. Puig, X. Obradors, J. Arbiol, H. C. Freyhardt, Adv. Function. Nanomat. 16 (2006) pp.1363–1372 M. J. Hÿtch, E. Snoeck, R. Kilaas, Ultramicroscopy 74, (1998) pp. 131–146
a) a)
b) b)
5 nm
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d)
Figure 1. (a) AFM image of CGO nano-huts oriented along two mutually perpendicular directions of the (001)-LAO surface; (b) HREM cross sectional image of a nano-hut viewed along the [1-10]CGO or [100]LAO directions. The interface shown in the image corresponds to a mismatch strain of ~30%. (c) εxx (d) εyy components of the strain versor within the island. The white numbers indicate the corresponding levels of strain.
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Diffraction contrast imaging and high resolution transmission electron microscopy of multiferroic thin films and heterostructures B.I. Birajdar, I. Vrejoiu, X.S. Gao, B.J. Rodriguez, M. Alexe, and D. Hesse Max Planck Institute of Microstructure Physics, Weinberg 2, D-06120 Halle (Saale), Germany [email protected] Keywords: transmission electron microscopy, multiferroics, epitaxial strain
Multiferroic materials exhibit simultaneous ferroelectric and ferromagnetic properties and are extensively studied because of their potential applications in multifunctional devices as well as for the understanding of the physics underlying multiferroicity [1]. The properties of these materials are influenced by their crystal structure and microstructural properties like strain, dislocations and domain boundaries, which can all be uniquely studied using transmission electron microscopy (TEM). Intrinsically multiferroic films and extrinsically multiferroic heterostructures and their components in the form of (1) BiFeO3 (BFO) epitaxial thin films (sample 1), (2) CoFe2O4 (CFO) epitaxial thin films (sample 2), (3) multilayers of (La,Sr)MnO3 (LSMO) and Pb(Zr,Ti)O3 (PZT) (sample 3), and (4) nanostructures of CFO (sample 4) (see Figure 1) were synthesised by pulsed laser deposition (PLD). Their microstructure was investigated using diffraction contrast imaging with a CM20T TEM operated at 200 kV, high resolution TEM [2] using a Jeol 4010 operated at 400 kV, and piezoresponse (PFM) and magnetic (MFM) force microscopies. Macroscopic ferroelectric and magnetic properties were measured by standard methods. The TEM bright field image of sample 1 in cross-section (Figure 1a) shows a 180 nm thick BFO film and a 50 nm thick SrRuO3 (SRO) bottom electrode on DyScO3 (DSO) (110) substrate. The films are single phase, show crystalline contrast and have sharp interfaces. The microstructure of the BFO films was extremely sensitive to the growth conditions (annealing temperature and oxygen pressure) and slight deviations from the optimum conditions led to secondary phases, in agreement with previous reports. The BFO film of sample 1 revealed a multidomain structure (Figure 1e) (cf. [3]), and similar domain structures were observed using PFM (Figure 1f). Diffraction patterns consisted of split spots confirming the formation of domains. Sample 1 has a remnant polarisation of about 55 μC/cm2 at room temperature and 1 kHz, similar to the value given in the recent literature reports [3], which is attributed to the high structural quality of these films. The TEM bright field image of sample 2 in cross section (Figure 1b) shows a 120 nm thick CFO film and a 20 nm thick SRO electrode on a STO substrate. The CFO film is epitaxial and single phase but has a high density of structural defects. In some films dislocations were observed at the SRO-CFO interface. (Investigations of the magnetic domain structure of the CFO films using plan-view TEM and MFM are under way).
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Indexing of [010] diffraction patterns showed a similar in-plane lattice parameter of the CFO film, the SRO film, and the STO substrate indicating strained epitaxial growth. The out-of-plane lattice parameter of CFO is larger than its bulk value leading to c/a ratio of 1.08.The CFO films appears to be in compressive stress. Sample 2 showed inplane and out-of-plane values of the remnant magnetization of 266 and 33 emu/cm3 respectively, which is comparable with recent reports in the literature [4]. Sample 3 is a graded heterostructure consisting of ferromagnetic LSMO films sandwiched between ferroelectric PZT films of increasing Zr content (Figure 1c) which exert in-plane tensile stress on the LSMO layers significantly altering their magnetic properties [5]. Sample 4 is an array of 35 nm large CFO islands (Figure 1d), synthesised using a deposition mask made from anodic aluminum oxide (AAO) [6]. Because of their reduced dimension, these CFO islands are expected to be under strain yielding significant changes in the ferromagnetic properties. The measurement of these properties is however challenging because of the complex design of nanostructured arrays. Quantitative analysis of the strain in the multilayers (sample 3) and nanostructures (sample 4) using HREM and geometric phase imaging (cf. [7]), as well as the application of other analytical techniques like EDX and EELS are under way. 1. 2. 3. 4. 5. 6. 7.
S.-W. Cheong and M. Mostovoy, Nature Materials 6 (2007) 13. Help by Dr. S. Senz and Dr. N. D. Zakharov in HREM is gratefully acknowledged. Y.-H. Chu et al., Adv. Mater. 18 (2006) 2307 W. Huang et al., Appl. Phys. Lett. 89 (2006) 262506 I. Vrejoiu et al., submitted to Appl. Phys. Lett. X. S. Gao et al., submitted to Int. Symp. on Integrated Ferroelectrics (ISIF), Singapore 2008. M.-W. Chu et al., Nature Materials 3 (2004) 87.
Figure 1. (a-d) TEM bright-field (cross section) images of samples 1-4, (e-f) TEM bright-field (plan-view) and in-plane PFM amplitude images of sample 1 showing a multidomain structure.
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Imaging of compositional defects at silicide-silicon interfaces using aberration corrected HAADF M. Falke1, U. Falke2, P. Wang2, A. Bleloch2 1. Institute of Physics, Chemnitz University of Technology, 09107 Chemnitz, Germany 2. SuperSTEM Laboratory, Daresbury, Cheshire, WA4 4AD, U.K. [email protected] Keywords: interfaces, HAADF, aberration correction,
Silicide-silicon interfaces were studied at the superSTEM Laboratory using aberration corrected dedicated STEM at 100kV. With suitable samples the resolution of about 0.1 nm allows compositional defects of atomic dimension to be visualised in high angle annular dark field (HAADF) images. Cubic metal silicides grown epitaxially on cubic (001) Si were chosen for these investigations, since, in case of bulk crystals on a [110] zone axis, only pure metal and pure silicon columns face the electron beam. At the interface however, strain and differences in chemical solubility can produce various, mostly periodically appearing one-dimensional defects. To evaluate the defect structure it is first necessary to understand the more or less undisturbed interface. We reported the interface structures found for cubic NiSi2 and CoSi2 on Si (001) [1,2]. Here we report some defects which disturb this interface structure and are usually found close to steps or local changes in composition. The periodic nature of the defects, shown here means that they could be described as compositional interface reconstructions. Sometimes the defect structure can be derived intuitively from the HAADF contrast, see e.g. Figure 1 showing missing metal columns which appear as dark triangles at the interface and Figure 2 showing that metal and silicon columns seem to exchange their places at a regular distance of a few unit cells. Image simulation and comparison with thickness and imaging data can reveal more about the real content of metal and silicon in the disturbed columns. Figure 3 shows the interface of a mixed NiSi/NiSi2 film on Silicon which was alloyed with 7% of platinum and annealed at 700 °C. Since Pt is not soluble in NiSi2 but easily soluble in NiSi the remaining Pt segregates at periodic sites at the NiSi2/Si interface. This effect is found close to nickel monosilicide crystallites, which contain Pt evenly distributed. Here again judging simply from the HAADF the bright spots arranged regularly at the interface are very likely to arise from Pt rich atomic columns. Image simulations support this and reveal more information about the Pt content. In summary we show how suitable and powerful aberration corrected HAADF imaging in a dedicated STEM is to study metal silicon combinations which are relevant to the semiconductor industry and which provide very interesting model interfaces. 1. 2. 3.
U. Falke, A. Bleloch, M. Falke, and S. Teichert, Phys. Rev. Lett. V92, N11 (2004) 116103. M. Falke, U. Falke, A. Bleloch et al., Appl. Phys. Lett. 86 (2005) 203103. We would like to acknowledge financial support from DAAD Project Number D/07/09995.
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Figure 1. Metal columns missing every fifth unit cell in HAADF contrast.
Figure 2. Periodically appearing compositional changes were observed close to interface steps in CoSi2 / Si samples. The inset shows an HAADF image simulation (E.J. Kirkland, Advanced Computing in Electron Microscopy, ISBN 0-306-45936-1).
Figure 3. HAADF image of the Si-A- to a NiSi2-B-type interface, containing Pt-rich columns.
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Characterization of nanometric oxide particles extracted from a steel surface onto a carbon replica P. Haghi-Ashtiani1, A. Ollivier2 and M.-L. Giorgi2 1. Laboratoire de Mécanique des Sols, Structures et Matériaux, CNRS UMR 8579, École Centrale Paris, Grande Voie des Vignes, 92295 Châtenay-Malabry, France 2. Laboratoire de Génie des Procédés et Matériaux, École Centrale Paris, Grande Voie des Vignes, 92295 Châtenay-Malabry, France [email protected] Keywords: replica TEM preparation, electron diffraction, selective oxidation
In steel industry, after cold rolling, steel sheets are annealed in an atmosphere of N2 and H2 (5 to 15 vol.%) in order to recrystallize the steel substrate. During annealing, the less noble alloying elements such as Si and Mn segregate to the surface where they form discrete oxide particles [1]. The main objective of the work presented here is to gain better understanding of the formation of these oxide particles. The annealing experiments were performed on a ternary alloy containing Fe, Si (0.094 wt.%) and Mn (0.492 wt.%), the principal impurities being C (less than 10 ppm), N, O and S (less than 5 ppm). Before the experiments, the substrates were mirrorpolished up to 1 µm with diamond paste. The specimens were annealed in the quartz chamber of an infrared radiation furnace (Ulvac Sinku-Riko) with a temperature profile in line with the industrial practice (Figure 1). The furnace gas atmosphere consists of a commercial high purity N2 and 5 vol.% H2 mixture containing a controlled partial pressure of water (12.8 Pa). After annealing, the laboratory steel surface is covered with small spherical particles (Figure 2a). The particles are extracted from the steel surface using the double replica methodology (Figure 2b): 1) in the first step, the particles are extracted onto a cellulose acetate polymer film, and 2) in the second step, the polymer film covered by the particles is coated with an approximately 20 nm-thick layer of carbon. This double replica is then inserted in a copper grid and the polymer film is dissolved carefully in acetone. The chemical nature of the particles is determined by electron diffraction in the transmission electron microscope (TEM JEOL 1200 EX) at 120 kV, using selected area diffraction and conventional bright field and dark field techniques. The experimental diffraction patterns are compared with theoretical ones in order to determine the chemical nature and orientation of the particle (Figure 3). The crystal structure of the possible oxides (lattice parameters and atom positions) is found in the Inorganic Crystal Structure Database (ICSD) [2] (Figure 3a). Based on the crystal parameters, theoretical electron diffraction patterns are calculated by means of the Java Electron Microscopy Simulation (JEMS) software [3] (Figure 3b). The particles present on the replica contain Mn2SiO4 (Figure 3c) and MnO. Conventional bright and dark field techniques reveal that the spherical particles can be
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composed of either several oxides or several grains of the same oxide with different orientations (Figure 4). This result illustrates the complexity of the oxides’ nucleation and growth within a particle. 1. 2. 3. 4.
A.R. Marder, Prog. Mater. Sci. 45 (2000), p.191. Fachinformationszentrum Karlsruhe, National Institute of Standards and Technology: Inorganic Crystal Structure Database ICDS, Karlsruhe, Germany, 2006. P. Stadelmann, Java Electron Microscopy Simulation (JEMS) software, CIME-EPFL, École Polytechnique Fédérale de Lausanne, Switzerland, 1999-2006. The authors are extremely grateful to Natural Resources Canada's CANMET-Materials Technology Laboratory for sharing their extraction replica methodology.
1000
T (°C)
800
Fig. 2a)
600
400
200
0 0
100
200
300
400
500
t (s)
Figure 1. Temperature profile
Fig. 3a)
Fig. 2b)
Figure 2. a) Specimen after annealing, b) TEM image of this specimen’s replica
Fig. 3b)
Fig. 3c)
Figure 3. a) Crystal of Mn2SiO4 in the [100] orientation, b) simulation for Mn2SiO4 in the [100] electron incidence, and c) experimental electron diffraction micrograph obtained
Fig. 4a)
Fig. 4b)
Fig. 4c)
Fig. 4d)
Figure 4. a) Bright and b), c), d) dark fields revealing different parts of the same particle
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TEM characterization of textured silicon heterojunction solar cells A. Hessler-Wyser1, C. Monachon1, S. Olibet2 and C. Ballif2 1. Interdisciplinary Centre for Electron Microscopy, Ecole Polytechnique Fédérale de Lausanne, 1015 Lausanne, Switzerland 2. PV-LAB, Institute of Microtechnology, University of Neuchâtel, 2000 Neuchâtel, Switzerland [email protected] Keywords: Heterojunction, thin films, HRTEM
Increasing the efficiency of crystalline silicon (c-Si) solar cells requires the reduction of both bulk and interface recombination. Even if bulk recombination is almost suppressed, the symmetry of the crystal lattice is disturbed at the surface and hence, due in particular to non-saturated bonds (dangling bonds), a large density of defects exists. Thus, in this case, the free-carrier lifetimes are no longer limited by the quality of the bulk c-Si, but by its surface. To keep the recombination losses at the c-Si surface at minimal levels, the surface must be electronically well passivated. An efficient way to obtain passivation is to use low temperature grown (typically 200°C) hydrogenated amorphous silicon (a-Si:H) [1]. In the case of photovoltaic applications, the passivation of both c-Si wafer surfaces (i.e., the emitter and the rear surface) is of crucial importance for good performances. In this work, the 3-7 nm a-Si:H based passivation layers of the device are grown by VHF-PECVD in a single chamber. The solar cells consist of a multilayered structure: Al back contact / DC-sputtered ITO (Indium Tin Oxyde) / n/i a-Si:H back surface field (BSF) / n-type c-Si substrate with a resistivity of 1-3 Ωcm / intrinsic a-Si:H / i/p a-Si:H emitter / and a front contact made out of DC-sputtered ITO via a shadow mask to define the cells (diameter = 4.5 mm). Such devices are called heterojunction (HJ) silicon solar cells. Surface recombination losses are a major concern for all c-Si solar cells. In particular, mastering of HJ emitter and back surface field formation on textured c-Si is crucial for high-performance HJ solar cells. High Resolution Transmission Electron Microscopy (HRTEM) is necessary to identify the key microstructural features of the aSi:H/c-Si interface, and TEM micrographs of HJ interfaces on flat and textured highperformance devices are shown for the first time and discussed with respect to the resulting solar cell electrical performances. TEM micrographs of our flat high-efficiency HJs show abrupt c-Si/a-Si:H/μc-Si:H interfaces for emitter and BSF formation. Whereas on the pyramidal facets of the textured substrate the growth is identical to the flat substrate interface, we observed unexpected epitaxial growth at the bottom of the pyramid valleys. We have identified these local epitaxial domains as an efficient surface recombination path, and by consequence, as responsible for the observed decreased VOC on textured HJ cells. When minimizing the density of epitaxial domains at the c-Si/a-Si:H interface by
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adapting the deposition conditions, a solar cell VOC of 660 mV is obtained. An additional modification of the textured c-Si surface morphology leads to VOCs as high as 700 mV. 1.
J. I. Pankove and M. L. Tarng, Appl. Phys. Lett. 34, (1979),156.
Figure 1. HR-TEM micrograph of a flat Si HJ solar cell emitter consisting (from left to right) of c-Si / intrinsic a-Si:H / p+ a-Si:H/μc-Si:H layer. (The complete device includes a stack intrinsic a-Si:H / n+ a-Si:H/μc-Si:H layer stack for the back surface.) This structure yields HJ devices with VOCs of up to 710 mV and efficiencies up to 19.1%.
Figure 2. TEM and HR-TEM micrographs of a textured Si HJ solar cell emitter consisting of c-Si / intrinsic a-Si:H / p+ a-Si:H/μc-Si:H layer. In comparison with Fig. 1, one can observe epitaxy of the a-Si:H-based passivation layer located at the bottom of the pyramid valleys. Minimization of the density of such local epitaxial domains leads to Si HJ solar cell VOCs over 700 mV. Therefore, we identify such local epitaxial domains as unsatisfactory passivated spots at the c-Si/a-Si:H interface.
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Investigation of the change in the microstructure of thin p-type Bi-Sb-Te thermoelectric films after heat treatment F. Heyroth1, M. Schade1, K. Rothe1, H.S. Leipner1 and M. Stordeur2 1. Interdisziplinäres Zentrum für Materialwissenschaften, Martin-Luther-Universität, 06099 Halle, Germany 2. angaris GmbH, Heinrich-Damerow-Str. 1, 06120 Halle, Germany [email protected] Keywords: thermoelectric film, LVSEM, EDX
Thermoelectric devices utilizing the Peltier and the Seebeck effect have been widely used for cooling and power generation applications in the last decades. Typically they are built from bulk materials. The energy conversion efficiency of such devices depends on the thermoelectric figure-of-merit (z⋅T) of the thermoelectric materials. In spite of numerous works to improve the z⋅T values of thermoelectric materials - which are dependent on the Seebeck coefficient S, the electrical conductivity σ, and the thermal conductivity κ (z = S²⋅σ / κ) - only a value around unity could be reached for bulk thermoelectric materials at room temperature. However, for low dimensional nanostructures, such as nanowires or superlattices, it has been predicted that the z⋅T values can be enhanced significantly compared to bulk materials due to quantum-confinement effects [1]. Beside this effect the progressive miniaturization requires thin thermoelectric films with Seebeck coefficients and electrical conductivities comparable to the bulk materials. Our films of the thermoelectric effective p-type compound semiconductor (Bi0.15Sb0.85)2Te3 were prepared by sputter-deposition. After heating of the as-deposited films and afterwards cooling, a distinct increase of the electrical conductivity σ can be observed. Measurements of the Seebeck and the Hall coefficient established that the increase of the electrical conductivity is not connected with an expected decrease of the Seebeck coefficient, because the charge carrier density is reduced but at the same time the hole mobility is increasing. The increase of the electrical conductivity at nearly unchanged Seebeck coefficient can be exploited for the enhancement of the film power factor (S²⋅σ). For the enlightenment of the changes in the film corresponding analytical investigations of the microstructure by XRD, EDX, and LVSEM were done. They show that, besides a grain growth in the polycrystalline films, a Te-rich phase appears after the heat treatment (cp. Figure 1). Also a generation of pores can be observed. The study of the crystalline nanostructure will be completed by EBSD and TEM investigations of the films. 1.
R. Venkatasubramanian, E. Siivola, T. Colpitts and B. O'Quinn, Nature 413 (2001) p. 597
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Figure 1. Left: Comparison between SEM micrographs of the cross section surface of the thermoelectric film (a) before and (b) after thermal treatment. Right: Corresponding EDX-line scans (integrated EDX mappings) along the cross section using a polished sample.
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EM study on forming Inorganic film with Periodically Organized Mesopores upon Polymer film Huang Wenqing1,2, Zhang Ying2, Yang Fang1, Zhaoxi2, Yang Wantai1 1. College of Material Science and Technology, Beijing University of Chemical Technology, Beijing 100029, China 2. Beijing Research Institute of Chemical Industry, SINOPEC. Beijing 100013, China [email protected] Keywords: Hybrid Thin Film, MCM-41, BOPP, PE
The potential use of nano-structured porous materials is currently recognized in advanced domains in materials chemistry, such as optical devices, photo-catalysis, catalysts, adsorption and sensors. The chemical nature and the morphology of the porous network can be accessed by XRD, IR, RAMAN, MAS-NMR, BET, SEM and HRTEM, which help us to make a better understanding of the nanocsale structure– property relationships and to control hybrid interfaces. [1] Here, we present a synthetic approach to produce Mesoporous materials containing designed nanoporosity on surface of polymer film after hydrophilic modification by confined photo-catalytic oxidation method. We use SEM to study the surface morphology and the thickness of the nanoporous film, and TEM to study the array of nanopores. In this work, a new synthesis method was studied to form a silica thin film with periodically organized nanopores upon polymer film. The UV irradiation system was used to make a modified polymer surface with high hydrophilicity[2]. Where, firstly, BOPP or PE films were cleaned with deionized water and acetone, and then predetermined amount of persulfate salt aqueous solution was deposited on the bottom film with a micro-syringe; a top layer transparent to UV light covered this solution and the drop of solution was spread into an even and very thin liquid layer. UV irradiation intensity is set at 8800uw/cm2, and irradiation time is controlled within 100s; then immersed the modified surface into water for 15hr. A MCM-41 sol-gel was spread to the treated surface by spin-coating method, and condensing reaction was conducted in 30min under acidic condition with triblock copolymer surfactant at room temperature. The surface morphology and the thickness of the inorganic film was showed respectively in Fig.1 and Fig.2. The thickness of the film range from 100nm to 2 micron can be controlled by spin coating speed and sol-gel concentration. TEM picture (Fig.3) showed the parallel pores and regular array of silica film after calcination. The diameter of the parallel pores is about 3 to 5 nm which is evidenced by Nitrogen Physisorption. 1. 2.
Clément Sanchez, Cédric Boissière, Chem.Mater. 20(2008), p.682-737. 2. P. Yang, J. Y. Deng, W. T. Yang, Polymer 44(2003), p.7157-7164.
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Figure 1. Surface morphology of inorganic fillm by SEM
Figure 2. Cross-section morphology of inorganic film upon polymer film by SEM
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Figure 3. TEM morphology of the powder after inorganic film calcination
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Nanointerface analysis of hard coatings deposited by IBAD D. Kakas1, B. Skoric1, A. Miletic1 and L. Kovacevic1 1. FTN – Centre for Surface Engineering and Nanotechnology, Trg D. Obradovica 6, 21000 Novi Sad, Serbia [email protected] Keywords: nanointerface, IBAD, cross section polisher
In this research hard TiN coatings were deposited onto previously carburized steel substrates (0.165%C, 0.2%Si, 1.2%Mn and 1%Cr). In order to sustain substrate properties obtained during heat treatment, hard coatings should be deposited at very low temperatures (up to 120°C). Special attention has been given to improving adhesion, which usually presents a problem when using low temperature deposition. Ion Beam Assisted Deposition offers a solution for this problem. However, IBAD alone is not enough since the main influence on adhesion has the interface between coating and substrate [1,2]. Ion Beam Mixing (IBM) has been used in order to acquire nanointerface between hard TiN coating and steel substrate. Using Ion Beam Mixing one can achieve intensive mixing of the atoms of a substrate with the atoms of a deposited coating. The thickness of the nano interlayer can easily be regulated by correction of parameters such as evaporation rate, energy of additional bombardment, atom to ion ration, incident angle etc. [3]. In this research nanointerface thickness was varied between 10 and 60 nm and its influence on coating adhesion was analysed. Adhesion of the coatings was characterized using scratch test equipment ST 2000. A diamond prism was used during the measurements. Damaged area was inspected using scanning force microscope. Micrograph of the scratch channel at the maximum load (100N) reveals that effects of delamination are not visible on the sides yet only at the end of the channel (Figure 1). Obtained results indicate very high adhesion. Our earlier investigations [2] show that other authors report much lower adhesion of coatings deposited at low temperatures. The presence of the nanointerface plays highly important role on excellent adhesion. This was confirmed using EDX analysis at the end of the scratch channel (Figure 1). Results of the EDX analysis are presented in the Table I. Spectrum 1 obtained at the place where coating delamination appeared shows that there is a 5.57 % of Ti in the surface of the base material. This means that coating delamination occurred across the nanointerface which was produced using IBM process. One can easily see that there are very small differences in composition in spectra 2, 3 and 4. It can be concluded that moving of diamond prism even at the load of 100 N didn’t change surface quality of TiN coating. In order to acquire the image of the nanointerface, samples were prepared using the Cross Section Polisher and inspected with HRSEM (JEOL JSM-7001F). The images were acquired at the JEOL (Europe) SAS Application Centre. Figure 2 shows the cross section of nanointerface.
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Figure 1. Micrograph of the hard the scratch channel is presented in the left, detail of delaminated area is presented in the right of the figure. Table I. Chemical composition at the end of the scratch channel (weight %) Spectrum
N
Ti
Cr
Mn
Fe
1 2 3 4
0.00 20.51 18.03 18.30
5.57 31.85 35.37 43.20
1.02 0.55 0.63 0.49
1.35 0.75 0.60 0.52
91.58 45.08 44.08 36.15
Figure 2. Micrograph of the hard TiN coating is presented in the left, high-resolution micrograph of a cross section of the nanointerface is presented in the right of the figure. Results of this investigation indicate a possibility of a new field of application of hard coatings such as deposition onto carburized steel substrates. Steel substrate provides high mechanical properties and TiN coating provides high wear resistance. 1. 2. 3. 4.
W. Ensinger, A. Schröer and G.K. Wolf, Surface & Coatings Technology 51, (1992), p. 217. D. Kakas, B. Skoric, T. Novakov, A. Miletic and L. Kovacevic, 10th International Conference on Tribology, Serbiatrib ’07 (2007), p. 71. B. Skoric, D. Kakas, M. Rakita, N. Bibic and D. Perusko, Vacuum 76, (2004), p. 169. This project was financially supported by the Ministry of Science and Environmental Protection of the Republic of Serbia.
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Transrotational crystals growing in amorphous Cu-Te film V.Yu. Kolosov, A.V. Kozhin, L.M. Veretennikov, C.L. Schwamm Physics Dept. Ural State Economic University, 8Marta 62, Ekaterinburg, 620219 Russia [email protected] Keywords: bend contours, internal crystal lattice bending, Cu-Te amorphous condensates
Crystallization in amorphous films often results in strong internal lattice bending: TEM reveals regular rotation of initial lattice orientation across the growing microcrystals [1]. According to electron diffraction data our deposits prepared by vacuum evaporation of Te layer upon Cu layer are initially amorphous in wide range of thickness and composition. Te crystals growing later in these films carry regular bend contour patterns on the TEM images useful [2] for careful analysis of lattice orientations and lattice bending. Crystallization was either spontaneous or initiated by the electron beam illumination. The crystals of hexagonal Te were formed more or less elongated (normally to [001]) with dominant orientations for the crystal centers <010> or <120>, Figure 1 (some grains with these orientations are marked). Such orientations correspond to most dense lattice planes lying almost parallel the surface (also parallel film plane). Bend contour technique was used to measure the magnitude of internal lattice bending (i.e. rotation of the lattice orientation) about two mutually perpendicular directions: around [001] and around its normal. Internal lattice bending around [001] is almost 4 times higher and is doubling (from 85 to 192 degrees per micrometer) upon film thickness reducing by half (from 40 nm to 22 nm), Figure 2. At the same time internal lattice bending around normal to [001] is increasing four times: from 12 to 50 and from 11 to 44 degrees per micrometer for orientations <010> and <120> correspondingly (in the central crystal areas where nucleation takes place). In addition the decrease of film thickness results in increase of aspect ratio. Anisotropy of internal lattice bending is caused by the character of chemical bonding (the covalent coupling of atoms in spirals lying along [001], and Van der Waals bonds between these spirals). Lattice orientation of the nucleus has less pronounced effect for a lattice bending (Figure 2). Qualitative estimates: the increase of Cu content rises internal lattice bending and reduces aspect ratio of the elongated grains, Figure 3. Often the total continuous rotation of the lattice orientation (corresponding to elastic bending) in the isolated crystals exceeds 180 degrees, Figure 4. Such large regular lattice misorientations (also published for other substances) unambiguously testify to the fact of internal (lattice) nature of the bending and cannot be caused by bending of a crystal as a whole body (for 180 degrees opposite corners of thin bent crystal should be nontransparent for electron beam). Thus the studied Te crystals can be associated with “transrotational” crystals [1]. 1. 2.
V.Yu. Kolosov and A. R. Tholen, Acta Mat. 48 (2000) p. 1829–1840. V.Yu. Kolosov, Proc. XII ICEM, ed. L. Peachey & D. Williams, (Seattle) 1 (1990) p. 574575.
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Figure 1. Micrographs of Te crystals with dominant orientations (marked for some grains by squares and circles) grown in areas of different film thickness (a – 38nm, b – 30nm, c – 22nm) and SAD (d)
Figure 2. Dependence of internal lattice bending in Te crystals upon the film thickness, measured for the grains from the upper Figure 1.
a
b
Figure 4. Te crystal with total lattice misorientation above 180˚ (from A to D). Internal lattice bending ~ 120˚/μm. ZAPs corresponding to <010> are marked with triangles. c
Figure 3. Te crystals in film areas with different Cu concentration (a – 22%, b – 20%, c – 18%). Aspect ratio goes down as the Cu content increases.
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TEM investigation of sputtered indium oxide layers on silicon substrate for gas sensors Th. Kups1, I. Hotovy2 and L. Spieß1 1. Institute of Micro- and Nanotechnologies, Materials of Electronics, Ilmenau University of Technology, P.O. Box 100565, 98684 Ilmenau, Germany 2. Department of Microelectronics, Faculty of Electrical Engineering and Information Technology, Slovak University of Technology, Ilkovičova 3, 812 19 Bratislava, Slovakia [email protected] Keywords: In2O3, PVD, HRTEM
Indium oxide (In2O3) has recently attracted much interest as a material for transparent electrodes in electronic devices such as liquid crystal displays [1], solar cells [2], as a barrier layer in tunnel junctions [3], as sensing material in gas sensors e.g. for ozone [4] and in nanowire technology [5]. Two crystal structures have been reported for In2O3: body-centered cubic (bcc, Ia3, a = 1.0118 nm) and rhombohedral (rh, R-3c, a = 0.5478 nm and c = 1.451 nm) [6]. In the case of bcc-In2O3, the physical and optical properties are well known [1-4], whereas for rh-In2O3, to the best of our knowledge, no information is available. Common deposition methods of In2O3 films and nanostructured layers are evaporation [7], magnetron sputtering [8] and Sol-Gel [9] processing. These deposition processes usually lead to the growth of In2O3 with a poly- or single crystalline bcc structure independent of the used substrate [10]. In this work magnetron sputtering technique will be used to form indium oxide layers and cross section high resolution transmission electron microscopy (HRTEM) for investigation of the crystal structure. Nanoscaled In2O3 films were deposited by dc reactive magnetron sputtering from a In target (76.2 mm in diameter, thickness of 3 mm and 99.99% pure) in a gas mixture of O2 and Ar. The distance between the target and substrate was approximately 75 mm. The apparatus was evacuated to a pressure below 5x10-4 Pa before deposition. A sputtering power of 75 W was used. Both the inert argon flow and reactive oxygen flow were controlled by mass flow controllers. The flow of oxygen in the reactive mixture O2-Ar was changed in the range of 40-80 sccm. The thicknesses of the resulting layers were in the range of 40-80 nm which were prepared onto unheated silicon substrates. Figure 1 shows the comparison of different layers sputtered at different growth parameters (Fig 1a, b: O2 flow 40 sccm compared to 80 sccm in Fig. 1c, d). The HRTEM images show different crystalline structures sizes and shape. While in the low oxygen flow case (Figs. 1a and b) the size of the crystallites is in the range of >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 [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 [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 [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) [email protected] 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 [email protected] 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
Series1
t = 48 nm
GaK / InK ratio * 0.351
Series2 Series3 Series4 Series5 Series6
400
100 nm 197 nm
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-
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 379–380, DOI: 10.1007/978-3-540-85226-1_190, © Springer-Verlag Berlin Heidelberg 2008
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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] 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.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 383–384, DOI: 10.1007/978-3-540-85226-1_192, © Springer-Verlag Berlin Heidelberg 2008
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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 [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.
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3. 4. 5.
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] 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 <111> 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 <111> 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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2.
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 <10 nm of the thinned foil for Si and W respectively. Redeposition of sputtered material will generally add more Ga on the specimens. This artefact can generally well be minimized with in-situ lift-out FIB preparation. Therefore, averaged over the specimen thickness only a minor Ga signal with weak materials dependence is to be expected in EDS analysis. Typical concentrations are reported to be in the range 0 at% [1] to < 2 at% [2]. Obviously these percentages will depend on the specimen thickness. These low Ga signals are indeed consistent with the results obtained in most of our specimens. However, a high gallium level is observed at the interface of Al layers as is illustrated on Figure 1 for an Al/SiO2/Si stack. Also in TiN/TiAl3/Al layer structures the Ga signal strongly peaks at the interfaces (Figure 2). In an early report [3] a high Ga signal observed in the TiN layer of a TiN/Al stack was attributed to the high sputtering yield of TiN. As the Ga ion beam basically mills parallel with the TEM foil, this explanation seems unlikely and more probably the lateral resolution during that measurement was insufficient to reveal that the Ga was actually peaked at the interface instead of in the TiN. The Ga accumulation at the interfaces can be understood by the rather low solubility of Ga in Al (<9at% at 26-29ºC). Grain boundary penetration of liquid Ga on Al bicrystals is extensively studied in literature [4-6] and shown to be a very fast process that can lead to embrittlement of the Al. The present observations show that this phenomenon can also occur at the interfaces of Al with other materials. 1. 2. 3.
C.R. Hutchinson, R.E. Hackenberg and G.J. Shiflet, Ultramicroscopy 94 (2003), p. 37. T. Ishitani, H. Koike, T. Yaguchi and T. Kamino, J. Vac. Sci. Technol. B 16 (1998), p. 1907. K. Park, Mat. Res. Soc. Symp. Proc. 199 (1990), p. 271.
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4. 5. 6.
R.C. Hugo and R.G. Hoagland, Scipta Materialia 41 (1999), p. 1341. W. Ludwig, E. Pereiro-Lopez and D. Bellet, Acta Materialia 53 (2005), p.151. E.E. Glickman, Z. Metallkd. 96 (2005), p. 1204.
150
Counts
Al
Si Al O Ga
100
SiO2
50
Si
a
0.0
b
50 nm
5
10 15 20 Position (nm)
25
30
Figure 1. HAADF-STEM image of the interface of a sputtered Al layer on SiO2/Si (a) and EDS line scan (b, top to bottom) across the interface showing the increase of the Ga signal near the interface of the Al and corresponding with the bright line on the STEM image. 100
SiGe
TiN
TiAl3
Al
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TiN
Counts
SiGe
Al
Ga
60
Ti Al
40
Si Ge
20
TiAl3
a
500 nm 500 500 nm nm
b
50
100
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Position (nm)
Figure 2. HAADF-STEM image of SiGe/Al contact (a) and EDS line scan (b, top to bottom) along the line indicated on the image showing the accumulation of Ga at the TiN/TiAl3 and TiAl3/Al interfaces (arrows).
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Precipitation in an Al-Mg-Ge Alloy R. Bjorge1, C.D. Marioara2, S.J. Andersen2 and R. Holmestad1 1. Norwegian University of Science and Technology, Department of Physics, 7491 Trondheim, Norway 2. SINTEF Materials and Chemistry, 7465 Trondheim, Norway [email protected] Keywords: precipitation, Al-Mg-Ge alloy, HRTEM, ADF-STEM
The precipitation in an Al-0.87 at% Mg-0.43 at% Ge alloy has been investigated using high-resolution transmission electron microscopy (HRTEM) and annular dark field scanning TEM (ADF-STEM). This work is a continuation of previous studies on AlMg-Si-(Cu) alloys, which are industrially relevant due to their superior mechanical properties such as high strength/weight ratio. The hardness increase is generated by the precipitation of nanometre-sized metastable phases from solid solution during artificial ageing. Most precipitates have had their crystal structures solved by our group using a combination of quantitative electron diffraction, first-principles methods and HRTEM [1-3]. It has been shown that all phases are structurally related through a near-hexagonal Si network having sub-cell dimensions a=b=0.405 nm, c=0.405 nm [3]. A possible way of improving the mechanical properties of Al-Mg-Si-(Cu) alloys is by replacing one or more solute elements with other ones having similar electrochemical properties. Total or partial replacement of Si by Ge is one such example. In bulk form Ge is similar in size to, and forms the same diamond structure as Si. An improvement in peak hardness and age-hardenability by replacing Si with Ge has previously been reported for an Al-1.0 wt% Mg2Ge alloy [4]. The higher atomic number of Ge with respect to Al and Mg makes Al-Mg-Ge alloys well suited for ADF-STEM studies. In the present work, hardness measurements indicate a moderate increase in strength and better stability during ageing for the Al-Mg-Ge alloy compared to a similar alloy containing Si instead of Ge. HRTEM and ADF-STEM images from early precipitates in the Al-Mg-Ge alloy show that Ge columns arrange in a near-hexagonal network with similar dimensions as the previously reported Si-network in Al-Mg-Si-(Cu) alloys. The Ge-network has directions parallel to <100>Al (Figure 1), producing precipitates that sometimes have plate-like morphologies. In most precipitates the atomic arrangement on the Ge-network was disordered, but a particle with a monoclinic unit cell was also found (Figure 2). This type of atomic arrangement has only been observed in the AlMg-Si-Cu alloys [5]. 1. 2. 3. 4. 5.
S. J. Andersen et al., Materials Science and Engineering A 390 (2005), p. 127. R. Vissers, et al., Acta Materialia 55 (2007), p. 3815. S. J. Andersen et al., Materials Science and Engineering A 444 (2007), p. 157. K. Matsuda, S. Ikeno and T. Munekata, Materials Science Forum 519-521 (2006), p. 221. C. D. Marioara et al., Philosophical Magazine 87 (2007), p. 3385.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 395–396, DOI: 10.1007/978-3-540-85226-1_198, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. Fourier-filtered ADF-STEM images of early-stage precipitates in the Al-MgGe alloy with the matrix in a <100>Al orientation. The image shows that the Ge columns are arranged in a near-hexagonal configuration with sub-cell dimensions a=b=0.405 nm, c=0.405 nm. The sub-cells are oriented parallel to <100>Al directions. This atomic arrangement is very similar to what has been found in the Al-Mg-Si-Cu system [5].
Figure 2. HRTEM image and corresponding FFT of a plate-like precipitate in the AlMg-Ge alloy viewed along <100>Al. Although the precipitate is partly disordered, a monoclinic unit cell of dimensions a=1.215 nm, b=1.032 nm, c=0.405 nm, γ=101.3° can be observed. The unit cell is also a super-cell in Al, as indicated.
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Interaction between dislocations and oxide precipitates in an aluminium containing ferritic stainless steel L. Boulanger, S. Poissonnet and F. Legendre Commissariat à l'Energie Atomique, DEN/DANS/DMN Service de Recherches de Métallurgie Physique Centre d'Etudes de Saclay, F91191 Gif-sur-Yvette, France [email protected] Keywords: precipitates, dislocations, strain contrast
Ferritic and martensitic steels strengthened by oxide dispersion (ODS) are studied at the nanometre scale because the interaction between dislocations and nanosized precipitates drives mechanical behaviour and evolution during neutron irradiation. Observation by Analytical Electron Microscopy (AEM) aside chemical information brings information about the coherency of precipitates, their strain field and the surrounding dislocations. We investigated the PM2000 alloy (Fe-19Cr-5.5Al-0.5Ti-0.5Y203 (wt.)) manufactured by Plansee. It contains a large range of precipitates which can be classified into two groups: respectively around 5 nm and 50 nm. The smaller contain Al-Y-O (1]. Our purpose is to show the contrast of nanoprecipitates under two beam imaging and their association to the dislocations. Two metallurgical states have been investigated, the asreceived state and an in-situ annealed obtained by heating inside the electron microscope at 930°C during 3 mn. In the as received material two dark lobes contrast “Figure 1”, so-called AshbyBrown contrast [2], is clearly visible on bright field images. Histograms sizes on out of contrast images give the same distribution. Among them, larger precipitates (> 10 nm) present parallel fine fringes attributed to moiré patterns “Figure 2”. An image compatible with a screw dislocation (b= ½[-1, -1, 1]) is visible on figure 1. Images realised with two perpendicular reflections of type (110) show a rotation of the lobes position “Figure 3a and 3b”. After in-situ annealing few evolutions is present: dislocations are still pinned on precipitates “Figure 4”. They show the very same strain contrast but sometime their square shape is better resolved. As conclusion, PM 2000 alloy contains very small precipitates (3 nm) clearly seen under two beams imaging. A dark contrast line is present that is aligned with the reflection vector according to a spherical strain. Above 7 nm moiré contrasts are present likely due to a relaxation of the stress. This structure is thermally stable up to 930°C. 1 2
M. Klimiankou, R. Lindau, A. Möslang and J. Schröder, Powder Met. 48(2005), p. 277. M. F. Ashby and L. M. Brown, Philos. Mag. 8(1963), p. 1649.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 397–398, DOI: 10.1007/978-3-540-85226-1_199, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. Two beams image
Figure 3a. First reflection: double lobes and dislocation line
Figure 2. Double lobes and moiré patterns
Figure 3b. Perpendicular reflection: rotation of lobes, the dislocation image vanishes
Figure 4. After annealing, the square shape becomes more visible
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Voids Associated with Nano-Particles of Tin in Aluminium Laure Bourgeois1,2, Matthew Weyland1 and Barry C. Muddle2 1. Monash Centre for Electron Microscopy, Monash University, VIC 3800, Australia 2. Centre of Excellence for Design in Light Metals, Monash University, VIC 3800, Australia Department of Materials Engineering, Monash University, VIC 3800, Australia [email protected] Keywords: vacancy clusters, aluminium alloy, transmission electron microscopy
Vacancies play a crucial role in many aluminium alloys: they mediate the diffusion of most solute atoms, relieve the compressive strain of precipitates, and must be present above a critical concentration for certain phases to nucleate [1]. Unfortunately vacancies and their association to solute atoms or precipitates can only be detected indirectly. In this work, transmission electron microscopy (TEM) is used to show that under certain circumstances three-dimensional objects containing a large amount of vacancies form in association with tin precipitates in aluminium. These objects bear some similarities to voids observed in other systems [2-3]. The examined samples were alloys of Al-0.01at%Sn and Al-1.7at%Cu-0.01at%Sn; these are model alloys for the study of micro-alloying additions (Sn in this case) on the precipitation behaviour of Cu-rich phases [4-5]. The samples were heat treated according to the usual regimes for precipitation-hardened aluminium alloys, with ageing temperatures ranging between 350 and 485°C. The TEM observations were performed on a JEOL 2011 operating at 200 kV. Figure 1(a) shows a bright-field image of a small cluster of tin particles observed in the Al-0.01Sn alloy. Each particle exhibits a single spheroidal region (see arrows) of contrast lighter than the dark tin particle. Figure 1(b) shows a lattice image of another such case. Energy dispersive x-ray (EDX) spectra collected with a focused 5 nm probe on different regions of the group shown in Figure 1(b) are presented in Figure 1(c). No elements other than aluminium or tin were found associated with the “bubble” contrast (the copper peak is an artefact). In particular, no enrichment in oxygen or other volatile element was detected. Examination of the diffraction contrast due to the “bubble” features over an extensive range of specimen holder tilts failed to detect any crystalline phases other than aluminium and copper. Furthermore, several examples were found of broken shells at the surface of tin particles, suggesting that the “bubbles” are empty. Further confirmation that these objects are either empty or at least contain significant amounts of vacancies was provided by the nature of the Fresnel fringe contrast at the matrix-void interface. Figure 2 shows the change in the Fresnel fringe at the matrix-void interface as a function of objective defocus. The fact that the white fringe faces the “bubble” region at the negative focus condition, and the matrix in the positive focus condition, is a clear indication of a lower projected charge density at the “bubble” region. These objects can therefore be regarded as voids. S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 399–400, DOI: 10.1007/978-3-540-85226-1_200, © Springer-Verlag Berlin Heidelberg 2008
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In the Al-0.01Sn alloy only a relatively small number of tin particles were observed associated with voids, and this only occurred at the high ageing temperature of 445°C, which is slightly less than the temperature where liquid tin dissolves into the matrix (460°C; note that the bulk melting point of tin is 232°C). However in the Cu-containing Al-1.7Cu-0.01Sn alloy, the voids were more common and persisted along with the tin particles well above the tin solvus in pure aluminium; the voids were often found at the triple junction of tin, aluminium, and a copper-rich grain. Sn is known to have a high vacancy binding energy. It is therefore not altogether surprising to find Sn particles associated with vacancy clusters. A possible mechanism might involve the nucleation of voids in the vacancy-supersaturated region surrounding a tin particle, the excess vacancies being liberated from solute tin atoms incorporated into the growing precipitate. The fact that the voids are often seen attached to a smaller grain on the surface of the tin particles suggests that they nucleate heterogeneously. 1. 2. 3. 4. 5.
K.C. Russell, Scripta Metallurgica 3 (1969), p. 313. C. Cawthorne and E.J. Fulton, Nature 216 (1967), p. 575. J.S. Williams, M.J. Conway, B.C. Williams and J. Wong-Leung 78 (2001), p. 2867. H.K. Hardy, J. Inst. Metals 80 (1951-52), p. 483. L. Bourgeois, J.F. Nie and B.C. Muddle, Philos. Mag. 85 (2005), p. 3487.
(a)
(b)
(c)
Figure 1. (a) Tin particles in Al-0.01Sn associated with spheroidal contrasts interpreted to be voids (see arrows); (b) lattice image of a tin-void group taken along the [100] zone axis of aluminium; (c) EDX spectra taken with a focused 5 nm probe for the different components of the group shown in (b).
Δf=-540nm
Δf=-72nm
Δf=+384nm
20 nm Figure 2. Part of a through-focal series, from underfocus (left) to overfocus (right). The Fresnel fringe contrast at the matrix-void interface as a function of objective defocus Δf confirms that the void region has a projected charge density that is less than the aluminium matrix and must therefore contain vacancies.
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Influence of thermal treatments in microstructure and recrystallization peak energy of P/M Al-Mg-X alloys S.J. Buso1,2, A. Almeida Filho1, I.M. Espósito1,2, J.R. Matos4, W.A. Monteiro2,3 1. Centro Universitário Sant’Anna – São Paulo - Brazil 2. Instituto de Pesquisas Energéticas e Nucleares IPEN/USP– São Paulo - Brazil 3. Universidade Presbiteriana Mackenzie – São Paulo - Brazil 4. Instituto de Química da Universidade de São Paulo IQ/USP– São Paulo – Brazil [email protected] Keywords: Aluminum Alloys, Thermal Analysis, Electron Microscopy, Recrystallization
Light materials have been studied thoroughly and used in components of pieces in the automobile, naval and aerospace industries in the last decade. Among those materials, the aluminium-magnesium alloys have special attention due, not only to the lightness of the material, but also to certain mechanical properties and reciclability[1,2]. In this work, samples of Al-2Mg-1Nb and Al-2Mg-0.6Zr alloys were produced, by powder metallurgy techniques, employing uniaxial hot compactation (400 MPa) and extrusion (60 MPa) both processes at 450oC. The product of these processes were bars of ∅ = 15 mm and length of 200 mm. The bars of the alloys were cold worked, without intermediate thermal treatments in a rolling mill being obtained reduction in area of 80%. Samples of both alloys in study were analyzed, without previous thermal treatments, at DSC in equipment SHIMADZU DSC 50 of the Laboratory of Thermal Analyses of the IQ/USP. From the same alloys lot, were prepared samples for TEM analysis in microscope JEM200C of the IPEN/USP. The figures 1.a and 1.b show the sequences of the analyses of DSC in samples of Al-2Mg-0.6Zr and Al-2Mg-1Nb alloys in temperature range between 400oC and 550oC, the curves shows the peak of the initial recrystallization energy for the alloys in study. The figures 2 to 3 show the microstructure observed by TEM of the samples in study. In the analyses of samples of Al-2Mg-0.6Zr and Al-2Mg-1Nb alloy, grains are observed with possible precipitates close to grain boundaries and inside the grains, in the samples without thermal treatment (figures 2.a and 3.a). After 60s of thermal treatment at 550oC (figures 2.b and 3.b) a distribution of precipitates and a strong interaction between dislocations and precipitates are observed. These processes are characteristic of the recovery and recrystallization. 1. 2.
Buso, S. J.; Monteiro, W. A. Microstructural characterisation of a P/M Al-Mg-Zr alloy after thermal treatmensts. Acta Microscópica, v. 9, p. 153-154, 2000. Buso, S. J., Monteiro, W. A.. Characterisation by TEM of a supersatured P/M Al-Mg-Zr alloy after thermal treatments In: International Conference on Processing & Manufacturing of Advanced Materials (Thermec’2003), 2003, Madrid, Proccedings of the International Conference on Processing & Manufacturing of Advanced Materials, 2003.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 401–402, DOI: 10.1007/978-3-540-85226-1_201, © Springer-Verlag Berlin Heidelberg 2008
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(a)
(b)
Figure 1. Sequences of the analyses of DSC in samples of: a)Al-2Mg-0.6Zr and
b) Al-2Mg-1Nb
(a)
(b)
Figure 2. Micrographies of samples of Al-2Mg-0.6Zr: a) without thermal treatment and
b) 60s, 550oC
(a)
(b)
Figure 3. Micrographies of samples of Al-2Mg-1Nb: a) without thermal treatment and b)
60s, 550oC
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Analytical TEM of Nb3Sn Multifilament Superconductor Wires M. Cantoni1, V. Abächerli2, D. Uglietti2, B. Seeber2 and R. Flükiger2 1. EPFL-CIME, Bat. MXC-134, Station 12, CH-1015 Lausanne 2. Group of Applied Physics, University of Geneva, CH-1211 Geneva [email protected] Keywords: TEM, EDX, FIB
Bronze processed Nb3Sn superconducting wires are today still considered as the best candidates for nuclear magnetic resonance (NMR) magnets operating in persistent mode at high magnetic fields. The metallurgical treatment of Nb3Sn bronze wires with additives like Ta and Ti is quite difficult to further improve because the correlation between the additive and its effect on the superconducting properties, which do not only depend on its concentration but also on the whole variety of fabrication parameters. The aim of this work is to clear up the different effects of the internal (Nb-Ta / Nb-Ti filaments) and external (Cu-Sn-Ti bronze) Ti doping method on the superconducting properties and to analyze their possible reasons [1]. Conventional techniques of specimen preparation did not lead to TEM-samples with the required large and thin areas for the analysis because of severe preferential thinning of the bronze matrix around the Nb3Sn filaments. The focused ion beam (FIB) was used to cut out lamellae of 20x5um size directly from a cable, to attach them to a TEM grid and to thin them further to electron transparency (Fig. 1a). The samples obtained this way have uniform thickness and are ideally suited for EDX-Analysis and elmental mapping. Already the raw intensitiy maps of the elemental distribution (Fig. 1b) allow an interpretation of concentration gradients. Transverse as well as longitudinal crosssections permitted the investigation of the grain structure and the elemental distribution parallel and perpendicular to the cable’s axis. Scanning Transmission Electron Microscopy (STEM) in combination with Energy Dispersive X-Ray (EDX) Microanalysis was used to analyze the chemical composition of the superconducting phase. A gradient of Sn was found across the individual filaments reducing this way the overall “useful” superconducting cross-section in high magnetic fields. EDX as well as energy filtered imaging (EFTEM) were used to investigate and trace the additives on a nanometer scale at the grain boundaries (Fig. 2). Cu could be found in almost all the grain boundaries where as the Ti was only present at the grain boundaries when the internal Ti doping method was applied. This finding allows to engineer grain boundaries and to control the superconducting properties of these multi-filament cables. 1. 2.
Abächerli V, Uglietti D, Seeber B, Flükiger R, Physica C 372-376 (2002), 1325-1328. We kindly acknowledge EPFL-CMI for the access to the FIB.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 403–404, DOI: 10.1007/978-3-540-85226-1_202, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. left) STEM dark field image: cross-section of a single superconductor filament. The inset shows the TEM lamella prepared by FIB. right) EDX element maps (raw intensity of selected x-ray peaks) of Nb and Sn. The remaining un-reacted Nb core and a gradient in the Sn concentration are clearly recognizable.
Figure 2. a) STEM dark field image (left) with EDX linescan (arrow). The linescan profiles (right) of Cu-L and Ti-K peaks show an increased concentration of Cu and Ti at the grain boundary.
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3D Reconstruction of Ni4Ti3 Precipitates in Ni-Ti by FIB/SEM Slice-and-View S. Cao, W. Tirry, W. Van Den Broek and D. Schryvers EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020, Antwerp, Belgium [email protected] Keywords: Ni4Ti3, 3D reconstruction, FIB/SEM
Ni4Ti3 precipitates with lenticular shape and rhombohedral atomic structure growing in the austenitic B2 matrix of binary Ni-rich Ni-Ti alloys upon proper annealing treatment have an important influence on the shape memory effect which originates from the martensitic transformation of B2－B19’ [1]. Figure 1 shows a typical 2D distribution of such precipitates as obtained by TEM under conventional imaging conditions. A slice-and-view procedure in a Dual-Beam FIB/SEM system has been applied to investigate the 3D distribution and morphology of these precipitates, in order to indicate their influence in the martensitic transformation. A bulk of material has been sliced away step by step using the Ga+ ion beam, and for each slice a 2D SE image has been taken. In this preliminary experimental setup the step size used is 100nm with a total of 55 steps, leading to an investigated volume of 136µm3. Due to the weak contrast of the precipitates, proper imaging conditions need to be selected to allow for semi-automated image treatment and an optimization procedure to find those has been followed. One of the SE images obtained from the Dual-Beam is shown is Figure 2(a)), proper alignment and image processing in MatLab© [2] and Amira© [3] are performed on the whole sequence to create a sequence of binary images (as one is shown in Figure 2(b)). A 3D reconstruction has been achieved via these 2D binary images (shown in Figure 3(c)), from which both qualitative and quantitative analysis can be performed. Finally, a volume ratio of 10.2% for the Ni4Ti3 precipitates could be calculated, summed over all variants, which yields a net composition of Ni50.27Ti49.73 for the matrix, leading to an increase of 125 degrees for the martensitic start temperature Ms. Also, the orientation between two precipitates could be measured as an average angle of 109.4° ± 1.1°, which confirms the relative orientation of the different variants of the precipitates between {111} family planes in theory (109.5°). The present work thus indicates that slice-andview is a valuable tool for the investigation of the 3D distribution and related parameters of technologically important precipitates in alloys such as Ni-Ti. In the near future, other quantitative measures on the distribution of the precipitates will be obtained. 1. 2. 3.
K. Otsuka, X. Ren, Progress in Materials Science 50 (2005), p. 511 Information on http://www.mathworks.com/products/matlab/
Information on http://www.amiravis.com/
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 405–406, DOI: 10.1007/978-3-540-85226-1_203, © Springer-Verlag Berlin Heidelberg 2008
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4.
We kindly acknowledge the help of MULTIMAT “Multi-scale modeling and characterisation for phase transformations in advanced materials”, a Marie Curie Research Training Network (MRTN-CT-2004-505226) of the European Commission for supporting this work.
Figure 1. Schematic drawing of the lens-shaped Ni4Ti3 in the two zones: (a) the [1 0 1]B2; (b) the [1 1 1]B2; (c) Typical TEM BF image showing four variants of Ni4Ti3 precipitates
Figure 2. One of the 2D cross-section images before (a) and after (b) background subtraction, and 3D reconstruction for all the Ni4Ti3 precipitates (c)
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Electron microscopy study of Mg78.5Pd21.5: a phase with nanothin 120° rotational twin domains W. Carrillo-Cabrera, J.P.A. Makongo, Yu. Prots, G. Kreiner Max-Planck-Institut für Chemische Physik fester Stoffe [email protected] Keywords: Mg78.5Pd21.5, rotational twinning, TEM study
Complex intermetallic alloy phases are intermetallic compounds with giant unit cells. They have a preference to accumulate at specific compositions. Therefore, it is difficult to obtain single phase materials and to determine their crystal structures. An example is “Mg4Pd”, first reported by Ferro [1]. No crystallographic data was known until Spiekermann [2] reported the large orthorhombic structure of Mg306Pd77 (Mg3.97Pd). In a study to find out the phase equilibria in the Mg-rich part of the Mg-Pd system [3] we additionally obtained three phases close to the composition Mg4Pd: Mg78.5Pd21.5 (Mg3.65Pd), Mg56.4Pd13.5 (Mg4.18Pd) and Mg57Pd13 (Mg4.38Pd). Our aim in this contribution is to demonstrate typical problems arising during the crystal structure determination of complex metallic alloy phases using Mg78.5Pd21.5 as an example. Its complex X-ray powder pattern (CuKα1) can be indexed on the basis of a hexagonal unit cell (ah = 28.046(8) Å, ch = 27.950(4) Å). A preliminary structure model was refined using single crystal X-ray data (Rint = 12.1%; wR2 = 29.3%, R1 = 16.4%). The calculated profile using this hexagonal model fits roughly the observed X-ray powder pattern (RB = 8.8%; wRp = 6.4%). However, this hexagonal model shows anomalous disordering features. Subsequent electron micrographs (Figures 1) clearly show that a “single crystal” specimen in reality consists of an intergrowth of narrow twinned planar domains parallel to the (001)h plane, which explains the hexagonal model inconsistencies. Further analysis demonstrated that the crystal structure is orthorhombic (a = 48.5729(9) Å, b = 28.0532(4) Å, c = 27.9859(5) Å, space group Fmmm) and the nanothin (5−50 nm) domains are 120° rotational twins oriented in three different directions perpendicular to the c-axis (Figure 2). FIB cuttings perpendicular and parallel to the c-axis were indispensable for this study. A more detailed study of the crystal structure and micro-structure of Mg78.5Pd21.5 (XRD, SAED, TEM, HRTEM) will be presented. 1. 2. 3. 4.
R. Ferro, J. Less-Common Met. 1 (1959), 424. S. Spiekermann, Doctoral Thesis, University Dortmund (1998). J. P. A. Makongo, Y. Prots, U. Burkhard, R. Niewa, C. Kudla. G. Kreiner, Phil. Mag. 86 (2006), 3. Part of this work has been done at the Special Electron Microscopy Laboratory at Triebenberg, Technical University of Dresden. We thank Prof. H. Lichte for his support.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 407–408, DOI: 10.1007/978-3-540-85226-1_204, © Springer-Verlag Berlin Heidelberg 2008
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[001]
001 1
2
2
3
1
2
3
Figure 1. TEM images of Mg78.5Pd21.5, showing nanothin domains (5−50 nm thick) perpendicular to the c axis. (left) overview, FIB cut parallel to the c axis. (right) Middle resolution TEM image of crushed particle, the layered domains are labelled as 1, 3 (viewed along [110]) or 2 (marked with arrows, viewed along [010]).
a 3* 3 2
220 000
1
b 3* [001]
Figure 2. Selected area diffraction patterns of Mg78.5Pd21.5 along [001]. FIB cut perpendicular to the c axis. (left) SAD of thick region with overlapped domains (1, 2, 3) and (right) SAD of one-domain region.
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Analysing small precipitates in a ferritic steel matrix A.J. Craven and M. MacKenzie Dept. Physics & Astronomy, University of Glasgow, Glasgow, G12 8QQ, UK [email protected] Keywords: electron energy loss spectroscopy, spectrum imaging, precipitates, steel
The strength of a microalloyed high strength low alloy steel is controlled by a number of factors including the grain size and the fine dispersion hardening precipitates. For sheet steel, the latter form after the final rolling pass, during and following the austenite to ferrite phase transformation. The majority of these particles are sub-10nm. Both 3 dimensional atom probe (3-DAP) and spectrum imaging (SI) using electron energy loss spectroscopy (EELS) have shown the peak in the size distribution in a Vbased steel is around 2nm [1]. As well as the size distribution, the composition of the particles is also of great interest. Earlier work using extraction replicas had shown that particles above 4nm in size were essentially nitrogen rich carbonitrides but there was some uncertainty as to whether a plasma process, used to thin the support film, was modifying the composition [2]. The more recent work, using both 3-DAP and EELS SI to analyse the particles within the ferritic matrix, confirmed that these results also applied to particles below 4nm in size [1]. 3-DAP had difficulty in quantifying the N content while confirming the low C content. In fact, the C seemed to form an “atmosphere” around the particles rather than being contained within them. SI was able to quantify the N but had difficulty determining the C content because of a perturbation of the background shape in front of the C K-edge. This perturbation is a major problem since its features are much bigger than the C K-edge intensity from a small particle. To avoid upsetting the objective lens with a large amount of ferromagnetic material, the specimen for EELS SI was made by focused ion beam lift out using 30keV Ga ions. The surface damage was removed using subsequent broad beam ion milling at 500eV. SI was carried out with a probe of half-angle 13mrad containing 0.19nA and using a collection half-angle of 27mrad and 1eV/channel. Figure 1 shows the residual shape left after spectra from the matrix had the background removed using a power law background model fitted in the region from 230 to 280eV. The spectra were recorded at different positions on the detector array and have been normalised to the first peak. Since the energies of the features are independent of the position on the detector, the perturbation arises from the matrix. It is ~2.5% of the background before the C K-edge. A method of modelling and accurately removing this perturbation is required. Since the particles are small and well separated, the surrounding matrix should provide a good model of the background. The particle positions are easily identified by mapping the intensity of the V L2,3 white lines. Thus the spectra in the pixels corresponding to the particles can be summed. If the size of the summed region over the particle is increased, the additional signal is purely matrix and this matrix contribution can be isolated by subtracting the sum obtained from the smaller region. Thus spectra from the particle plus matrix and the matrix alone can be obtained.
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Since the perturbation comes from the matrix, which is essentially Fe, the two spectra are normalised by the counts in the corresponding regions of an Fe L2,3 map. Figure 2 shows two such normalised spectra while Figure 3 is their difference. This difference is just the spectrum from the particle itself and is only ~1% of the total signal cf 2.5% for the background perturbation. While the difference is noisy, it is clear there is little intensity from the C K-edge but significant intensity from the N K-, V L2,3- and Mn L2,3-edges. The intensity from the Fe L2,3-edges has gone but there is some differentiation of the white lines due to jitter on the accelerating voltage. Similarly the O K-edge signal from the surface oxidation has been removed. Figure 4 shows the sum of the spectra extracted from 5 particles and compares it to spectra from stoichiometric VC and VN standards. These have been scaled to have the same intensities over the energy range 510 to 620eV. It is clear that the C content in the particles is low or nonexistent while the N content is close to that of stoichiometric VN. Thus EELS SI has the ability to analyse particles in the sub-4nm size range within a steel matrix but care must be taken with the processing of the data. 1. 2.
A.J. Craven, M. MacKenzie, A. Cerezo, T. Godfrey and P.H. Clifton, Mat. Sci. Tech. in press. J.A. Wilson, A.J. Craven, Y. Li and T.N. Baker, Mat. Sci. Tech. 23 (2007) 519. 1E +6
3 E +6 En e rg y Lo s s ( e V) 230
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Figure 2. Spectrum from particle plus matrix (thin black) and matrix (thick grey)
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Figure 3. The difference between the spectra in Figure 4. Background subtracted sum of spectra Figure 2. 1% of the signal is from the particle. from four particles compared to VC and VN.
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Failure analysis of first stage land-based gas turbine blades F. Delabrouille1, F. Arnoldi1, L. Legras1 and C. Cossange1 1. EDF R&D Les renardières, department MMC, 77818 Moret sur Loing, France [email protected] Keywords: superalloy, crack, Inconel 738LC
The failure of a Gas Turbine (GT) first stage blade was investigated by visual inspection, scanning electron microscopy, electron probe microanalyser and transmission electron microscopy. The objective of this studies is to characterize and identify the microstructural changes involved in this failure. In order to carry out examinations and analyses, sections of the airfoil were cut using spark machining at half-height of blade. The blade was characterized using a LEO SUPRA 35 SEM. A TECNAI F20 operating at 200 kV is used for all TEM imaging. The macrostructure examinations show a lot of grains in the blade wall. The grains have a lengthened shape near the surface and is relatively equiaxe inside. This blade was made of the widely used Ni-base superalloy Inconel 738LC. Inconel 738LC have a typical superalloy microstructure, it’s a multiphase microstructure consisting of γ matrix, bimodal γ’ precipitates, γ-γ’ eutectics, interdendritic MC and M23C6 carbides. Two main zones of damage were observed : two cracks on the airfoil tip and one crack on pressure side near platform cooling hole. The examinations was carried out on one crack on the airfoil tip. The cross-section examinations show that the crack are initiated at the interface between thermal barrier coating and bond coat located under spallation zones. The crack growth was mainly interdendritic, with some transgranular zones (Figure 1). When the crack is interdendritic the matrix is not deformed, the deformation is only observed in the zones where the crack is transgranular (Figure 2). The metallographic examinations were completed by detailed observations of the crack tip using transmission electron microscopy. The crack tip is located in a MC carbide (Figure 3) . The crack changes of direction at the carbide / metal interface and restart perpendicularly in the primary carbide. The crack propagates along crystallographic steps about 10 nanometres to form an arch of circle up to the crack tip. The width of the crack is about 200 nm at the level of the observed zone. The EDX analysis didn’t show significant enrichment in oxygen on the crack lips and no variation of the contents of elements of the alloy. The mixed character of propagation, the absence of creep cavities and the presence of fine striations in initiation zone let suppose that thermo mechanical fatigue is the main mechanism responsible of cracking. The absence of oxide at the crack tip implies that the crack propagation is due to thermomechanical loading at low temperature. The current work indicates that the mechanism seems to be the out of phase thermomechanical fatigue. However, it’s not possible to eliminate totally a secondary influence of creep and high cycle fatigue.
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Figure 1. Cross-section examination of crack at airfoil tip (SEM, backscatter detector).
Figure 2. Orientation map near the crack tip at the airfoil tip.
Alloy
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Carbide Figure 3. Bright field image of the crack tip at the level of a primary carbide.
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TEM investigation of microstructures in low-hysteresis Ti50Ni50-xPdx alloys with special lattice parameters R. Delville1, D. Schryvers1, Z. Zhang2, S. Kasinathan2, R.D. James2 1. Electron Microscopy for Materials Science (EMAT), University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium 2. Department of Aerospace Engineering and Mechanics, University of Minnesota, Minneapolis, Minnesota 55455, USA [email protected] Keywords: shape-memory alloy, low hysteresis, twinning
NiTi is one of the most popular shape-memory alloys (SMA) in medical applications because of its biocompatibility and its remarkable properties that allow recoverable mechanical energy to be stored in a compact delivery system. Still, its undesirable fatigue properties exemplified by the occurrence of medical-device fracture, along with large temperature/stress hysteresis and the narrow temperature range of operation translate to a tight margin of error for engineering design of the devices. A new theory on the origin of reversibility of phase transformations predicts the microstructures and explains the fundamental cause of large transformation hysteresis commonly shown by SMAs [1]. It predicts that the hysteresis can be drastically minimized by improving the geometric compatibility of the martensite and the austenite. This has been demonstrated [2] with alloys with special lattice parameters with a middle eigenvalue of their transformation matrix satisfying λ2=1 [2,3,4]. Their hysteresis decreases as the middle eigenvalue gets closer to one. The present work focuses on the TEM investigation of ternary Ti50Ni50-xPdx alloys in which different amounts of Pd substitution on Ni lead to special ratios between the austenite and martensite lattice parameters. As the compatibility between the two phases increases, a change in the microstructure is observed. Away from the compatibility condition with λ2>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 <2-11> 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<1, type I/II twins are forbidden and only compound twins allowed. As the content of Pd is decreased toward the compatibility condition, detwinning of the martensite is observed. When the compatibility condition is satisfied (Ti50Ni39Pd11) the microstructure contains large plates of untwined martensite (figure 1). Internal microtwins have completely disappeared. Since the distortion is minimal, twins do not need to be fine. Random twinning is observed, contrasting with ordered stacks of twins observed with alloys away from the compatibility condition. Twinning seems to occur
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along boundaries of martensite plates or as a result of nucleation variation. In addition to the (1-11) type I and <2-11> type II twins, (011) compound twins are also observed. This can be accounted to the fact that the sample sits at the limit condition λ2=1, thus allowing the three type of twins to coexist in the sample due to local lattice parameters variation. An ongoing study of local variation of composition (EDS, EELS) or lattice parameters (CBED) is attempting to relate them to the type of twins. In-situ study of the growth of martensite inside the austenite is made possible at the Ti50Ni40Pd10 composition which contains nucleated martensite inside the austenite at room-temperature. Preliminary results show exact austenite-martensite habit planes with no traces of twinning. It agrees with the fact that Ti50Ni40Pd10 is also very close to satisfying the compatibility condition. In-situ cooling will show how such a twinless martensite can grow along the habit plane and how the final martensite microstructure arises. 1. 2. 3. 4. 5.
Ball, J. M., James R.D., Phil. Trans. R. Soc. Lond., (1992) A 338-339. J. Cui et al., Nature Materials 5 (2006) 286-290. K.A. Bywater, J.W. Christian, Phil. Mag. 25-26 (1972) 1249-1273. W.J. Moberly, J.L. Proft, T.W. Duerig, R. Sinclair, Mat. Sci. For., 56-58 (1990) 605-610. We kindly acknowledge the support from the Marie Curie Research Training Network MULTIMAT “Multi-scale modelling and characterization for phase transformations in advanced material” (MRTN-CT-2004-505226)
Figure 1. Change of microstructure with composition is shown in the bright field images (a)-(c). Fig. (a) shows internally twinned martensite plates, (b) shows the detwinning of the plates and (c) a twinless plate when the compatibility condition is satisfied. SAD patterns (d,e) correspond to plates A and B in (a), respectively, showing a (1-11) type I twinning. SAD pattern (f) was taken over the twin boundary C and shows a <2-11> type II twin.
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Evidence of silica layer at the interface between ferrite and the chromium oxide scale in oxidized Fe-Cr-Si alloys G. Bamba , P. Donnadieu, Y. Wouters and A. Galerie SIMAP, INPGrenoble-CNRS-UJF, BP 75, 38402 Saint Martin d'Hères – France [email protected] Keywords: TEM, stainless steels, oxidation, silica, adhesion
The kinetic law of oxidation of Fe-Cr-Si ferritic steels strongly depends on the Si content [1]. In addition, the chromia scale formed during oxidation exhibits strong spallation on cooling function of oxide thickness but no spallation for Si-free alloy, up to a scale thickness of 2.7 μm. Such observations are usually attributed to the presence of silica formed at the steel-chromia interface, modifying both oxidation kinetics and scale-metal adhesion [2]. Though a continuous silica layer is frequently invoked as an explanation for enhanced spallation, mainly in austenitic stainless steels, few morphological details of this layer have been produced. A series of Fe-15Cr-Si ferritic stainless steels, containing 0, 0.5 and 1 wt.% Si, were oxidized in different atmospheres between 850°C and 950°C. In all cases, SEM investigations were unable to evidence details at the interface between the chromia scale and ferrite. TEM has been then carried out on two Fe-15Cr-Si alloys which have been oxidized at 850°C under two different atmospheres : 150 mbar H2O for 72h (sample 1), 150mbar O2 for 24 h (sample 2). Cross section samples suitable for TEM observations of the oxide layers have been prepared by the FIB technique. In sample 1 (Fig. 1), a continuous layer of silica is detected, as attested by the diffraction pattern taken in this area (Fig. 1b) and the composition measured by EDX (Fig. 1c). This silica layer is about 0.5 μm thick and displays a regular shape (Fig. 1d). In some places, the silica layer grows more deeply inside the ferrite grains. The chromia scale is formed by small grains (about 0.1 μm). It is worth noting that the silica layer is separated from the chromia scale by small ferrite grains. Figure 2 shows the FIB section and the interfacial area in sample 2. A silica layer is observed, but it is discontinuous and much thinner (about 50 nm) than in sample 1. The chromia scale is still formed by a fine grain microstructure. As illustrated here, a TEM study of FIB cross section can easily evidence the silica layer formed during oxidation even for very thin layers. Preparing the section with the FIB technique was essential because of the brittleness of the interface which prevents from using classical methods. TEM gives also a measurement of the silica layer as well as the chromia scale which are usually estimated indirectly by thermogravimetric analysis. This TEM investigation points out the different behaviour of the alloys in terms of silica layer formation such information should be particularly important for the understanding of the oxidation mechanism and of the spallation effect. 1. 2.
G. Bamba, Y. Wouters, A. Galerie, F. Charlot, A. Dellali, Acta Mater. 54 (2006) p. 3917. H.E. Evans, D.A. Hilton, R.A. Holm, S.J. Webster Oxid Met 19 (1983) p. 1.
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Chromia scale
Silica layer
Silica layer ferrite
Figure 1. TEM image of the interfacial area observed in the FIB cross section from sample 1 (Fig. 1a), electron diffraction pattern (Fig. 1b) and EDX spectra (Fig. 1c) taken in white layer area in Fig. 1a. High magnification image of the silica layer (Fig. 1d) and of the chromia scale (Fig. 1e) Chromia scale Chromia scale
Silica layer Silica layer
Figure 2. TEM image of the interfacial area in sample 2 (Fig. 2a). details of the interface between the chromia scale and the ferrite (Fig. 2b). The silica layer is ~ 50 nm thick and does not form a continuous film.
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Applying a classical 2 beam diffraction contrast method for measuring nanoprecipitate misfit L. Lae1, P. Donnadieu2 1. Société BASSETTI, 91bis avenue du Général Mangin, 38100 Grenoble France 2. SIMAP, INPGrenoble-CNRS-UJF, BP 75, 38402 Saint Martin d'Hères [email protected] Keywords: misfit, precipitate, 2-beam conventional TEM
Hardening precipitation remains of high interest for improving the mechanical properties of light alloys for structural applications. This precipitation occurs in the earliest stages of decomposition of a supersaturated matrix. The nanoscale precipitates which are often metastable are stabilized by elastic and interfacial energy. Hence particular attention is given to the misfit between the precipitates and the matrix. The precipitates being out of equilibrium, the misfit can be known only by direct measurement, basically by TEM because of the precipitate size (~ 5 to 20 nm). TEM provides numerous methods to measurement local deformation like CBED or HRTEM combined to the Geometric Phase Analysis. However because of the high particle density typical of hardening precipitation, CBED cannot be used as well as HRTEM when the precipitates are spherical. Ashby and Brown [1] have earlier established a method to measure the coherency strain for spherical precipitates using conventional TEM. Still the method has been rarely used quantitatively. As most of the CCD cameras are now able to record high contrast images, it seems interesting to revisit the possibility of the “Ashby and Brown” method. We have applied it to Al3(Sc,Zr) precipitates within the Aluminum matrix [2]. The composition of the precipitates is close to Al3Sc, though with a small amount of Zr . Figure 1a show the simulation of the BF image of a precipitate of radius r = 10 nm. The simulation is based on the Howie-Whelan equations under 2 beam conditions (g = (220)). The description of the displacement field is obtained using Eshelby description. Anomalous absorption and image forces were taken in account, the sample thickness is assumed to be 100 nm. As illustrated by Figure 1b, the simulated image intensity profile is very sensitive to the misfit value. Hence the Ashby and Brown method stands as a quite sensitive and appropriate method to measure the coherency misfit. Besides comparison of intensity profile can be simply used for determining the misfit. Figure 2 compares the experimental image and the experimental intensity profile to the one simulated for a misfit of 0.3 %. The method has been repeated on a limited number of precipitates. For precipitate with radius 9-10 nm, the misfit is 0.3 % while for larger precipitates (radius 14 nm), the misfit is 0.8 %. In any case the measured misfits are significantly smaller than 1.4 %, the value for bulk Al3Sc. The increase of misfit with the precipitate size is consistent with the high value for the equilibrium phase. However even the larger precipitates are far from the equilibrium value. Such difference points out the interest for direct
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measurement on nanoprecipitates in view of a modelling taking in account the elastic energy. It should be emphasized that owing to the digitization facilities given by CCD cameras, the classical method of Ashby and Brown seems to be very relevant and quite simple to carry out for quantitative study of coherency strains of nanoparticles. 1. 2. 3.
M. F. Ashby and L. M. Brown Phil. Mag. 8 (1963) p.1083 A. Deschamps, L. Lae and P. Guyot, Acta Materialia, 55 (2007)p. 2775 Part of this work has been funded by the joint research program Precipitation between Pechiney, Usinor, CNRS, and CEA
b
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Figure 1. Simulated Dark Field image of a coherent precipitate (radius 10 nm, misfit 0.2 %) (Fig. 1a) and the intensity profile of simulated images for misfit : 0.2 %, 0.5 % and 1% (Fig. 1b)
a
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Figure 2. Experimental BF image under 2 beam condition (Fig. 2a) Comparison of experimental and simulated intensity profile (misfit : 0.3 % ; precipitate size 9 nm, measured with DF images made on a superlattice spot).
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Microstructure and interface composition of γ-phase in Co38Ni33Al29 shape memory alloy R. Espinoza1, B. Bartova1, D. Schryvers1, S. Ignacova2 and P. Sittner2 1. EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium 2. Institute of Physics ASCR v.v.i., Na Slovance 2, Prague 8, CZ-182 21, Czech Republic [email protected] Keywords: Shape memory alloys, interface, FIB
Ferromagnetic shape memory alloys (FSMAs) are being intensively studied because of their potential applications as smart materials. In these materials, the martensitic transformation processes can be triggered by applying a magnetic field, in addition to the conventional mechanism in SMAs. Recently, Co–Ni–Al alloys have attracted increasing interest as new FSMAs [1-4]. Due to their physical and mechanical properties they are also attractive as moderate high-temperature SMAs. The Co-Ni-Al system undergoes a martensitic transformation from β-phase (cubic ordered B2-type) austenite to L10 (tetragonal) martensite in a temperature range between -180 and 120°C [5]. However, the alloy is extremely hard and brittle in the polycrystalline state, and it was found that the presence of the secondary γ-phase significantly improves the poor ductility of Co–Ni–Al polycrystals [2, 4, 6]. Nevertheless, the exact role of the γ-phase in the martensitic behavior is not clear, e.g. by inducing nucleation points for the transformation, affecting composition, etc. On the other hand, the γ-phase is not homogeneously distributed at a small enough scale (Figure 1) [7], that easily allows the presence of an interface between phases in a regular TEM sample prepared by electropolishing. Because of this, an adequate study of the interface requires the use of an advanced preparation technique like Focused Ion Beam (FIB) to obtain samples for TEM. The objective of this work is to characterize the microstructure and composition of a Co38Ni33Al29 alloy in the interface between γ phase and the β- or martensite phases. After the identification of the γ-phase in a polished surface, the FIB technique allowed us to prepare samples at the desired position (Figure 2). The microstructure of the γphase was not affected during the sample preparation, showing a L12 ordered structure. The results obtained by EDX showed a local variation of the composition near the interface with the β-phase (Figure 3), which means a depletion of the Co content of around 6% (see Table I). 1. 2.
H.E. Karaca, I. Karaman, Y.I. Chumlyakov, D.C. Lagoudas, X. Zhang, Scripta Mater. 51 (2004), p. 261. R.F. Hamilton, H. Sehitoglu, C. Efstathiou, H.J. Maier, Y. Chumlyakov, X.Y. Zhang, Scripta Mater. 53 (2005), p. 131.
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3. 4. 5. 6. 7. 8.
Z.H. Liu, X.F. Dai, Z.Y. Zhu, H.N. Hu, J.L. Chen, G.D. Liu, G.H. Wu, J. Phys. D Appl. Phys. 37 (2004), p. 2643. R. Kainuma, M. Ise, C.-C. Jia, H. Ohtani, K. Ishida, Intermetallics 4 (1996) S151. K. Oikawa, L. Wulff, T. Iijima, F. Gejima, T. Ohmori, A. Fujita, K. Fukamishi, R. Kainuma, K. Ishida, Appl. Phys. Lett. 79(20) (2001), p. 3290. P.J. Brown, K. Ishida, R. Kainuma, T. Kanomata, K-U. Neumann, K. Oikawa, B. Ouladdiaf, K.R.A. Ziebeck, J. Phys. Condens. Mat. 17 (2005), p. 1301. B. Bartova, D. Schryvers, Z.Q. Yang, S. Ignacova, P. Sittner, Scripta Metall. 57(1) (2007), p. 37. Acknowledgements: R. Espinoza would like to thank to the Belgian Science Policy for the financial support. B. Bartova and S. Ignacova gratefully acknowledge financial support from the MULTIMAT Marie Curie Research Training Network (MRTNCT-2004-505226).
Figure 1. Microstructure of Co38Ni33Al29 alloy consisting of β-phase matrix with uniformly distributed γ-phase.
Figure 2. SEM image showing martensite and γ-phase at room temperature in the Co38Ni33Al29 alloy. Table I. Composition obtained by EDX of the points showed in Figure 3. Point 1 2 3 4 5 6 7
Figure 3. CoNiAl FIB TEM sample. The lighter zone in the middle (delimited by lines) corresponds to the γ-phase. The numerated points indicate the positions of EDX analysis.
Co 35.023 31.127 53.250 52.977 55.258 35.239 37.007
Ni 31.072 30.831 27.946 29.849 28.707 31.146 32.010
Al 33.682 37.326 18.239 16.868 15.683 33.465 30.454
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Microstructural characterization of the aluminum alloy 6063 after work hardening treatments I.M. Espósito1,2, S.J. Buso1,2, W.A. Monteiro2,3 1. Centro Universitário Sant’Anna – São Paulo - Brazil 2. Instituto de Pesquisas Energéticas e Nucleares IPEN/USP– São Paulo - Brazil 3. Universidade Presbiteriana Mackenzie – São Paulo - Brazil [email protected] Keywords: Aluminum Alloys, Thermal Treatments, Electron Microscopy, Hardness
The aluminum 6063 alloy possesses a great industrial interest, presenting characteristics that justify its frequent use, when compared to the other aluminum alloys: the precipitation hardening and high cold work capacity. These alloys present high ductility, that allows their use in operations with high deformation degrees, as the cold work.[1,2] The aluminum alloys of the 6xxx series are used in a wide variety of applications, from the use in architecture profiles, as it is the case of the alloys more diluted and consequently of smaller mechanical resistance, as to 6063 and 6060, until the named alloys of structural application with larger tenors of alloying elements and higher hardness, with great application potential in the automotive industry, among other. [3] For this present work, samples of the commercial aluminum 6063 alloy bar (φ = 8 mm) was cold worked (90% area reduction) and then submitted to isothermal treatments at 735 K at 0, 60, 600 and 6000 s. Following the processes of analysis, samples for TEM were obtained, being the process of preparation the usual route. From the same lot a series of samples were prepared for hardness analysis. The observation of microstructure was made in a TEM JEOL JEM 200 C, available in CCTM of IPEN/USP and the hardness analysis in MICROMET Buehler from the Mechanical Engineering Department from USP, with a charge of 100 g. Figure 1 shows the evolution of microstructure of the alloy during the thermal treatments: 1a). a net of dislocations due to the high plastic deformation in sample without thermal treatment, 1b) a contour of grains showing a net of dislocations inside one of them due to the presence of precipitates Mg2Si for sample with 60 s of treatment, 1c) dislocations in an contour of grain, for sample with 6000 s of treatment. Figure 2 shows the hardness Vickers curve obtained for the samples in study. The analysis shows a decrease in hardness from 0 to 600 s of time of treatment, due to the recovery process in the final stage and simultaneously the recrystallization process. From the 600 s of treatment the hardness increases slightly due, mainly, to the presence of precipitates Mg2Si. 1. 2.
Espósio, I.M.; Buso, S.J.; Monteiro, W.A. Hardness analyses in an 6063 alloy after cold work and thermal treatments. In: International Materials Symposium, Materials 2005, Aveiro, Portugal, 2005. METALS HANDBOOK, ASM, 9th printing, v.14 Metals Park, Ohio, 1988, p.317-319
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Espósio, I.M.; Buso, S.J.; Monteiro, W.A. Análise da liga comercial 6063 após tratamentos termomecânicos. In: Congresso Brasileiro de Ciência e Engenharia de Materiais, CBCIMAT 2006, Foz do Iguaçu, Brazil, 2006.
(b)
(a)
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Figure 1. TEM of samples of aluminum 6063 alloy 90% cold worked: 1a). a net of dislocations due to the high plastic deformation in sample without thermal treatment, 1b) a contour of grains showing a net of dislocations inside one of them due to the presence of precipitates Mg2Si for sample with 60 s of treatment, 1c) dislocations in an contour of grain, for sample with 6000 s of treatment.
Figure 2. Sequences of the analyses of DSC in samples of Al-2Mg-1Nb
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Microstructure- mechanical property relationships in dual phase automotive strip steels V. Tzormpatzdi, and G. Fourlaris Laboratory of Physical Metallurgy, National Technical University of Athens, 9 Heroon Polytechniou St., 15780, Athens, Greece. [email protected] Keywords: dual phase, microstructure, mechanical properties, martensite
Dual phase (DP) and Transformation Induced Plasticity (TRIP) steels belong to the family of advanced high strength strip steels (AHSS) [1]. These steel grades are developed to permit the down-gauging of the body-in-white vehicle components, while providing high strength, suitable σ0.2% / tensile strength ratios, excellent formability and crash performance. The increased use of modern automobiles leads to the production of safe, fuel efficient and environmentally friendly vehicles. The objective of the present study is to assess the effect of variable intercritical annealing sequences to the microstructure and the attendant mechanical properties of two different dual phase strip steel grades. Intercritical annealing experiments were carried out for the DP800 and DP600 grades within the temperature range of 750 to 875oC, for time intervals ranging from 180 to 450s followed by water quenching. The microstructure was investigated using scanning electron microscopy (SEM) and X-ray diffraction. The mechanical properties were assessesd via Vickers hardness testing, while selective samples were tensile tested using standardized samples under static deformation conditions. The microstructure of the as received DP 800 steel is presented in Fig. 1, while that of the as received TRIP 600 is given in Fig.2. The experimental results showed that following intercritical annealing of the DP800 at 775oC for 180s the mechanical properties of the modified DP800 steel were enhanced (i.e. higher elongation values coupled with an improved σ 0.2%/tensile ratio of 0.48, while maintaining a UTS value of 800MPa-Fig.5). This has been attributed to the microstructural features of the modified DP800 product (Fig.3), i.e. refined martensite island dispersion at ferrite/ferrite grain boundaries, while maintaining a martensite volume fraction of 39%. Similar attempts of enhancing the mechanical properties of DP600 were met with less success, since a significant increase of the tensile strength of intercritically annealed DP600 samples was measured, for the intercritical annealing temperatures and cooling regimes employed (Fig. 4). The present study has demonstrated that under carefully controlled processing conditions it is entirely feasible to obtain modified microstructures in DP800 grades with a desirable mix in terms of the martensite to ferrite ratio, and with suitable martensite particle dispersions characteristics, leading to the production of enhanced DP800 strip steel products. Such a goal has not been accomplished, as yet, for the DP600 grade.
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1. Senuma T. ISIJ International, 41 (2001) 520-532 2. Rocha R.O., Melo T.M.F., Pereloma E.V.,Santos D.B.Mat. Sci. and Eng.A, 391(2005) 296-304
Figure 1. SEM micrograph of DP800 strip steel in the as received condition
Figure 2. SEM micrograph of DP600 strip steel in the as received condition
Figure 3. DP800 intercritically annealed at Figure 4. DP600 intercritically annealed at 775oC for 180sec, water quenched 775oC for 180sec, water quenched
Figure 5. Curve presenting the σ 0.2% / UTS versus UTS
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Electron diffraction analysis of nanocrystalline Fe-Al C. Gammer, C. Mangler, C. Rentenberger and H.P. Karnthaler Physics of Nanostructured Materials, Faculty of Physics, University of Vienna, Boltzmanngasse 5, 1090 Vienna, Austria [email protected] Keywords: Electron diffraction, peak broadening, nanocrystalline materials
Finite grain sizes lead to a broadening of diffraction peaks. For defect free grains the median grain size (Dmed) can be determined from the half width at half height (∆g) of a diffraction peak using the Scherrer equation (Dmed=0.443·(∆g)-1) [1]. In the present work this equation is applied for using TEM methods to determine the size of grains, the size of coherently scattering domains (CSD) and that of ordered domains. Disordered nanocrystalline Fe-45at.%Al was produced by high pressure torsion (HPT) from B2 ordered material [2]. Samples annealed to different temperatures were investigated. Their diffraction profiles were deduced by ring integration of the TEM diffraction images and the peak broadening was determined by fitting Voigt peakfunctions. Images were taken using a TEM operating at 200kV. The instrumental peak broadening can be deduced from the width of diffraction spots of a single crystal. To compensate for this, the results were deconvoluted. Figure 1 shows a bright-field image and the diffraction pattern of a sample that is disordered and nanocrystalline after HPT; the grain boundaries are not well defined. Annealing the sample to 370°C (cf. Fig. 2) leads to defect recovery and to the restoration of the B2 structure. The latter correlates with the appearance of the (100) superlattice reflection (cf. Fig. 2b). From the width of the different reflections in the diffraction profiles (cf. Fig. 3) median domain sizes were calculated (cf. Fig. 4). The deduced values using fundamental reflections (110) and (200) reflect the CSD because of the high density of defects within the grains. With increasing annealing temperature the CSD size increases. Using TEM dark-field images (with fundamental reflections) the grain size at 370°C was found to be log-normal distributed with a median of 38±4 nm, that is higher than the CSD size as not all the defects are recovered. The ordered domain size estimated from the (100) peak width approaches the CSD size and is in good accordance with results from TEM dark-field images using a (100) reflection. From the present results it can be concluded that in nanocrystalline materials the Scherrer equation is a very powerful tool for using TEM methods to estimate median grain and domain sizes using both fundamental and superlattice reflections. 1. 2. 3.
F. Fultz and J.M. Howe, “Transmission Electron Microscopy and Diffractometry of Materials” (Springer Verlag, Berlin Heidelberg) (2008), p.425 C. Mangler, C. Rentenberger, I. Humer, L. Reichhart and H.P. Karnthaler, Microscopy and Microanalysis 13 (2007), p.298 We acknowledge support by the research project “Bulk Nanostructured Materials” of the research focus “Materials Science” of the University of Vienna and by the Austrian research fund FWF.
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Figure 1. Fe-45at.%Al deformed by HPT with a shear strain of 10 000%. (a) TEM bright-field image. (b) corresponding diffraction pattern.
Figure 2. Fe-45at.%Al deformed by HPT (10 000%) and annealed at 370°C. (a) TEM bright-field image. (b) corresponding diffraction pattern.
Figure 3. Diffraction profile of Fe45at.%Al deformed by HPT (10 000%) and annealed at 170°C. Superlattice reflections are indicated in bold.
Figure 4. Variation of the median coherently scattering domain sizes compared to the ordered domain sizes.
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Dual Beam and TEM characterisation of deformation structures in fatigued austenitic stainless steel A. Garcia1,2, L. Legras1, M. Akamatsu1 and Y. Bréchet2 1. EDF R&D, Les Renardières, Département MMC, 77818 Moret sur Loing, France 2. SIMAP, INPGrenoble, CNRS UJF, Domaine Universitaire, BP 75, 38402 Saint Martin d’Hères cedex, France [email protected] Keywords: austenitic stainless steel, fatigue, dislocations structures, TEM, Dual Beam.
Thermal fatigue cracking is a damaging mode of austenitic stainless steels, particularly in the mixing zone of the Residual Heat Removal system (RHR system) of Pressure Water Reactor (PWR). The aim of this research project is to study and compare the deformation microstructures, at crack-tips and in the bulk, observed on a pulled out structural component made of AISI 304L stainless steel and those developed during cyclic deformation, on low cycle fatigue standard specimens. Such a comparison of the deformation microstructures could give key indicators of the real magnitude of the mechanical solicitation undergone by the pulled out structural component. First, this study is devoted to identify the microstructures developed in the bulk of fatigued AISI 304L specimens. For this purpose, total-axial-strain-controlled low cycle fatigue tests, performed in air and PWR water, at 150°C and 300°C, with 0,10% and 0,50% total strain amplitudes, are used. The microstructural characterization is carried out by Transmission Electron Microscopy (TEM), on a FEI Tecnai G2 F20 operating at 200kV. At the lowest strain amplitude, for each temperatures and environments, planar dislocation arrays and stacking faults on {111} easy glide planes prevail. Whereas, at the highest strain amplitude, the superposition of planar slip on {111} planes with the ladder-like or wall and channel structure of persistent slip bands (PSB), presented in Figure 1a, and cells predominate. Among the three variable parameters of the fatigue tests, the strain amplitude is the most influent one on the observed deformation microstructure. These structures are known to be characteristic of austenitic stainless steel with intermediate stacking fault energy (SFE) [1-4]. Misorientation between cells is also measured in TEM, thanks to the ACT facility developed by Dingley and Wright [5], which uses reconstructed diffraction patterns from conical dark fields. Besides, deformation microtwins, presented in Figure 1b, are detected in PWR water at 300°C for both strain amplitudes. The twinning plane is the (111) plane which is determined by a combined analysis of the SAD pattern [6] with the transformation matrix for deformation twinning calculated by Mahajan [7]. Second, the RHR system’s microstructure is compared to the results previously presented in this paper. Planar dislocation arrangements and microtwins along {111} planes are observed in the pulled out specimen, which is similar to the microstructures detected in standard fatigued specimen cycled at low strain amplitude and high temperature.
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Last, the microstructures observed along cracks of the RHR component are analysed by TEM, using cross section samples prepared using GATAN PIPS, and by an FEI Helios Nanolab Dual Beam, equipped with an HKL Nordlys CCD camera and with 3D imaging and EBSD softwares (AutoSliceAndView and AutoReveal). It allows to prepare thin foils in very precise zones along the crack, at its vicinity and at the cracktip. Cross section samples, localised on the structural component’s crack, points out the presence of oxide at the crack-tip and of nano-steps on the crack edges. Moreover, the low dislocation density observed at the crack tip indicates that the residual plastic stresses and the plastic deformation are rather low in this area. Works are carried out by Dual Beam in order to build a 3D representation of the dislocations structures, from series of Electron Channeling Contrast Images (ECCI), using the Forescatter detector (FSD) located at the bottom of the EBSD detector. 1 2. 3. 4. 5. 6. 7.
F. Ackermann, L.P. Kubin, J. Lepinoux, and H. Mughrabi, Acta Metall. 32 (1984), p. 715. M. Gerland, J. Mendez, P. Violan and B. Ait Saadi, Mater. Sci. Eng. A118 (1989), p. 83. C. Laird, Z. Wang, B.T. Ma and H.F. Chai, Mater. Sci. Eng. A113 (1989), p. 245. K. Obrtlík, T. Kruml and J. Polák, Mater. Sci. Eng. A187 (1994), p. 1. S.I. Wright and D.J. Dingley, Materials Science Forum 273-275 (1998), p. 209. T.H. Lee, C.S. Oh, S.J. Kim and S. Takaki, Acta Mater. 55 (2007), p. 3649. S. Mahajan, Metallography 4 (1971), p. 43.
Figure 1. TEM micrographs of AISI 304L steel cycled with Δεt/2 = 0,50 % (a) coexistence of cells, PSB and planar slip (b) microtwins in PWR water at 300°C. 80 70
Ni Fe Cr
60
g = [ 1 10]
Ni Fe Cr
50 40 30 20 10 0 0
10
20
30
40
50
100 nm
Figure 2. TEM micrographs of a secondary crack in the RHR system (a) propagation with nano-steps (b) Presence of oxide at the crack-tip revealed by an EDX profile.
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Microstructural characterisation of steel heat-treated by the novel quenching and partitioning process K. He1, D.V. Edmonds1, J.G. Speer2, D.K. Matlock2 and F.C. Rizzo3 1. Institute for Materials Research, School of Process, Environmental and Materials Engineering, University of Leeds, Leeds LS2 9JT, United Kingdom 2. Advanced Steel Processing and Products Research Centre, Colorado School of Mines, Golden, CO 80401, USA 3. Department of Materials Science and Metallurgy, Pontifícia Universidade CatólicaRio de Janeiro; RJ 22453-900, Brazil [email protected] Keywords: steel, quenching and partitioning, microstructure
The Quenching and Partitioning (Q&P) process, a new development in steel heat treatment, is designed to create mixed martensite/retained austenite steel microstructures [1,2]. It involves partial transformation of austenite to martensite by interrupted quenching to a temperature (QT) between the martensite start and finish temperatures, followed by annealing at a partitioning temperature (PT) designed to enrich the remaining austenite with carbon, thereby stabilizing it against further transformation to martensite upon subsequent cooling to room temperature. Cementite precipitation in martensite is suppressed by alloying with Si or Al, thereby increasing the time for carbon to escape supersaturation by diffusion into adjacent untransformed austenite. This philosophy is similar to that used to produce carbide-free bainitic structures [3], of current interest in the development of TRIP-assisted automotive sheet steels [4]. Potential advantages of Q&P are: high strength from martensite fraction; exceptionally high carbon concentration of austenite fraction, which is decoupled from austenite decomposition as compared with carbon partitioning during the bainite transformation; fast processing limited only by the rapid diffusion of carbon. Microstructural evolution during Q&P processing of medium-carbon austenite (steel composition: 0.6C-0.95Mn-1.96Si-wt%) for various quenching and partitioning temperatures and partitioning times was examined using light optical and transmission electron microscopy. Comparison was made between the microstructural characteristics of fully austenitised Q&P treated samples and ones subjected to bainitic treatments. Figures 1 and 2 illustrate microstructures after Q&P treatment compared with conventional bainitic treatment. Both Q&P samples, partitioned at 400 and 500°C, possess uniform fine microstructure after short annealing times (Figs 1a and 1b). However, characteristically, the bainite transformation was incomplete at 400°C after shorter annealing times, only partially decomposing the austenite (Fig 1c), and required a much longer time to develop similarly uniform microstructures of carbide-free bainite (Fig 1d). Examination by TEM similarly reveals a finer, more uniform Q&P structure of martensite laths interwoven with retained austenite (Fig 2a and b) produced at the
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shorter annealing times compared with the bainite structure of bainitic ferrite interwoven with retained austenite (Fig 2c) achieved in the same steel composition. These results indicate that Q&P processing can produce refined homogeneous microstructures with substantial fractions of stabilized retained austenite, and generally within a time scale likely to be much shorter than that of a conventional bainite treatment, which may be of importance to industrial steel heat treatment procedures. 1. 2. 3. 4.
J.G. Speer et al., Acta Mater. 51 (2003), p. 2611. D.V. Edmonds et al., Mater. Sci. Eng. A438-440 (2006), p. 25. H. K. D. H. Bhadeshia and D. V. Edmonds, Metall. Trans. 10A (1979), p. 895. D. K. Matlock and J. G. Speer, in “Proc. 3rd Int. Conf. Structural Steels”, ed. H.C. Lee, (The Korean Institute of Metals and Materials, Seoul, Korea) (2006), p. 774.
(a)
(c)
(b)
(d)
(a)
(b)
(c)
Figure1. Light optical micrographs: (a) QT=190ºC and PT=500ºC for 10s; (b) QT=190ºC and PT=400ºC for 120s; (c) bainitic treatment 400ºC for 120s; (d) bainitic treatment 400ºC for 900s.
Figure2. TEM micrographs: (a) QT=190ºC and PT=500ºC for 10s (bright-field and centred dark-field image using (220)γ); (b) QT=190ºC and PT=400ºC for 120s (centred dark-field image using (002)γ); (c) bainitic treatment 400ºC for 120s.
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Martensite tempering behaviour relevant to the quenching and partitioning process K. He1, D.V. Edmonds1, J.G. Speer2, D.K. Matlock2 and F.C. Rizzo3 1. Institute for Materials Research, School of Process, Environmental and Materials Engineering, University of Leeds, Leeds LS2 9JT, United Kingdom 2. Advanced Steel Processing and Products Research Centre, Colorado School of Mines, Golden, CO 80401, USA 3. Department of Materials Science and Metallurgy, Pontifícia Universidade CatólicaRio de Janeiro; RJ 22453-900, Brazil [email protected] Keywords: steel, tempering, carbide
The Quenching and Partitioning (Q&P) process is a steel heat treatment to create microstructures with retained austenite [1,2]. It involves quenching austenite to between the martensite start and finish temperatures, followed by a partitioning treatment to enrich the untransformed austenite with carbon, thereby stabilizing it to room temperature. Carbon partitioning to austenite, however, is competitive with carbide precipitation [2] and so Si and/or Al alloying is suggested and employed in order to suppress cementite formation to protect carbon. However, previous studies show that alloying is ineffective against transitional epsilon carbide precipitation [2]. Consequently, it is important to understand the effects of Si and Al on tempering behaviour relevant to a Q&P treatment of steel. Microanalytical electron microscopy has thus been used to determine the compositions of epsilon carbide and cementite and the epsilon to cementite transition during tempering of quenched Si- and Al-containing high-purity Fe-C alloys. In Fe-0.6C-1.97Si, tempered in the range 200-500°C, transitional epsilon carbide and equilibrium cementite were both observed (Figure 1), but within different time frames, as indicated by the schematic transition line in Figure 2. Epsilon carbide forms first but is succeeded by cementite after increasing times at decreasing temperatures, e.g. 5, 10 and 70 min at 500ºC, 450ºC and 400ºC, respectively. STEM-EDX analysis of epsilon carbides and adjacent ferrite using a 1nm probe showed no composition difference between them i.e. the transitional carbide possessed a Si concentration similar to that of the surrounding matrix. In contrast, TEM-EDX analysis of cementite precipitates detected negligible Si. These recent analyses are confirmation of early reported studies [3-5] and reinforce the generally accepted argument that Si influences the formation of cementite because its low solubility requires rejection during carbide formation [3]. Knowledge of the temperature-time transition line would be useful for optimizing the partitioning condition for Q&P treatment of steel. For the two partitioning conditions indicated in Figure 2, even though both lie within the epsilon carbide region, T1t1 is favoured over T2t2, because at higher tempering temperature the driving force for epsilon carbide formation is reduced and this condition is closer to the transition from S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 431–432, DOI: 10.1007/978-3-540-85226-1_216, © Springer-Verlag Berlin Heidelberg 2008
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epsilon to the more stable cementite phase. In addition, epsilon carbide formation must compete with the escape of carbon from martensite to austenite in a Q&P treated microstructure, calculated to be extremely rapid [6] (especially at higher temperatures, although probably hindered by trapping at sites within the martensite). With decreasing carbon supersaturation, the driving force for carbide precipitation is also further reduced. In Fe-0.6C-2.41Al, TEM observations showed that Al addition gave similar results to Si in terms of carbide composition and precipitation behaviour. As for Si alloying, Al did not suppress formation of epsilon carbide. At 450ºC the epsilon carbide to cementite transition was observed at around 10min, similar to that for Si alloying. (a)
(101)ε+(101)
(b)
(010)ε
(c)
(d)
(-11-1)ε
Figure 1. Fe-0.6C-1.97Si: (a) STEM bright-field image of epsilon carbides, 300°C for 15min; (b) bright-field and centred dark-field images using (-1,1,-1)ε reflection, 410°C for 15min; (c) corresponding electron diffraction pattern; (d) centred dark-field image using (312) cementite reflection, 460°C for 15min.
Figure 2. Schematic TemperatureTime diagram of the epsilon carbide to cementite transition.
1. 2. 3. 4. 5. 6. 7.
J.G. Speer et al., Acta Mater. 51 (2003), p. 2611. D.V. Edmonds et al., Mater. Sci. Eng. A438-440 (2006), p. 25. W.S. Owen, Trans ASM 46 (1954), p. 812. J. Gordine and I. Codd, J. Iron Steel Inst. 207 (1969), p. 461. W. C. Leslie and G.C. Rauch, Metall. Trans. 9A (1978), p. 343. F.C. Rizzo et al., in “Solid-Solid Phase Transformations in Inorganic Materials”, ed. J.M. Howe et al., (TMS, The Minerals, Metals & Materials Society, Warrendale, PA, USA) (2005), p. 535. Acknowledgement to the UK North-West Fine-Probe Microanalytical STEM Facility.
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Chemical and structural analysis of NiAl-Al2O3 interface by FETEM and STEM W. Hu1, T. Weirich2 and G. Gottstein1 1. Institut of Physical Metallurgy and Metal Physics, RWTH Aachen University 2. Central Facility for Electron Microscopy, RWTH Aachen University [email protected] Keywords: interface structure, chemistry, FETEM, STEM
In continuous ceramic fiber reinforced intermetallic matrix composites (IMC) the interfacial shear strength plays an essential role for the load transfer from the matrix to the fiber and determines the strength and ductility of the IMCs. The interfacial strength is controlled by the interface structure and chemistry which developed during the IMCfabrication process. The present investigation focused on the evolution of interface structure and chemistry of NiAl coated Sapphire-fibers during composite fabrication (fiber coating, diffusion bonding and embedded casting). The interface structure of the as-coated fiber is shown by a bright field EFTEM micrograph in Figure 1a. It reveals a layer of several 10nm thickness between fiber and NiAl. HRTEM images demonstrate that this interface layer has an amorphous structure (see Figures 1b and c) and consisted of oxygen, aluminium and nickel, as shown in the element mapping taken by EFTEM (Figures 1d-f). The composition of the layer was semi-quantitatively analysed with STEM/EDX. It contained about 60 at.% oxygen, 20~30 at.% Al and 10~20 at.% Ni. Thus, this amorphous interface layer was identified as a compound (Al,Ni)-oxide. This amorphous interface formed by incorporation of Al and Ni with the remaining moisture of the residual gas in the reaction chamber during the vapor phase growth in high vacuum during the initial stage of the PVD-process, and caused a reactive thin amorphous oxide film on the fiber surface. After diffusion bonding (1300 °C/40 MPa/1h) the amorphous interface layer still existed between fiber and matrix (Figure 2a), but many nanocrystalline particles precipitated either in the amorphous layer or at the boundary between the layer and NiAl as well as the fiber (Figure 2b). Element maps obtained by EFTEM reveal several regions in the layer with enrichment of Al or Ni. The regions with high Al concentration corresponded to those of high oxygen content (Figures 2d and e). The areas of high Niconcentration are well separated from those of high O/Al-concentration (Figure 2c). We assume that these Al/O- or Ni-enriched areas correspond to those with nanocrystalline precipitates which were observed by HRTEM (Figure 2b). Furthermore, there existed also a thin continuous sublayer with a thickness of about 2~3 nm between the NiAl and the interface layer that contains large amounts of Al and oxygen (Figures 2d and e) and thus, was identified as Al-oxide. The interface structure and chemistry of NiAl-composites after casting at about 1700 °C are shown in Figure 3. There is still a continuous layer at the interface, merely the layer has a thickness of about 2~3 nm (Figure 3a). HRTEM reveals that this thin layer
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consists of two different sublayers. An element line profile across the interface from positions A to D, taken by a STEM/EDXS-analysis suggests that the interface thin layer is composed of two sublayers: one is the Ni-rich sublayer at position C (thickness<2 nm) adjacent to the NiAl-matrix. This sublayer is likely a nickel-rich nickel aluminide with a Ni/Al-ratio ≈ 1.17. The other sublayer consists of Al-oxide and is located at position B (thickness<2 nm) in vicinity of the Al2O3-fiber (Figure 3b). These two sublayers of nickel aluminide and Al-oxide shown in the element line profile apparently correspond to the ones displayed in Figure 3a. Formation of an Al-oxide sublayer at the boundary with NiAl (Figure 2) and with nickel-rich nickel aluminide (Figure 3) led to a NiAl/Al2O3-bonding that caused a high interfacial shears strength due to a large work of adhesion between Al NiAl -O and NiAlAl 2 O 3 and consequently a high strength of NiAl composites. a
c
Al K-edge
e
NiAl NiAl
Al 2 O 3
Amorph. layer
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d
O K-edge
f
Amorph. layer
Al 2O 3
2 nm
N i L 2 3 -e d g e
c
a a
O K -e d g e
e
Figure 2. Interface structure of asdiffusion bonded NiAl composite (a) bright field EFTEM image shows the interface structure, (b) HRTEM image shows the fine structure of the interface layer. (ce) element mapping in the area of Figure 2a, (f) results of semiquantitative STEM/EDX-analysis at the locations A-E in Figure 2a.
A l2O 3 A B N iA l C
E D
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Figure 3. Interface structure and chemistry of an as-cast composite. (a) a bright field EFTEM image reveals the interface structure, (b) element profile across the interface along the white line in Fig. 3a.
Ni
600
400
Al
Counts
a
f S T E M /E D X -a n a lys is (a t.% ):
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d
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200
D
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Figure 1. Interface structure and chemistry of as-coated fiber. (a) bright field EFTEM image shows the interface structure, (b-c) HRTEM image reveals the fine structure of the interface, (d-f) EFTEM element mapping in the interface area shown in Figure 1a.
0
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TEM investigations of aluminum precipitate in eutectic Si of A356 based alloys Z.H. Jia1, L. Arnberg1, P. Åsholt2, B. Barlas3, T. Iveland2 1. Department of Materials Science and Engineering, Norwegian University of Science and Technology, Alfred Getz vei 2b, N-7491 Trondheim, Norway 2. Hydro Aluminium R&D Sunndal, N-6601 Sunndalsøra, Norway 3. Montupet SA, 60181 Nogent sur Oise Cedex, France [email protected] Keywords: precipitate; aluminum alloy; electron microscopy
Solidification of hypoeutectic type A356 alloys start with growth of primary aluminum (α-phase), followed by Al-Si eutectic [1]. Precipitation of Si, together with other added elements from aluminum has been extensively studied [2, 3]. However, little detailed characterization has been done on the eutectic Si phase. Recently, nanoscale precipitates have been observed in the Si phase [4, 5] and have been confirmed to be aluminum [6], which exhibit three preferential orientation relations with the Si matrix [001]Si║[111]Al; [111]Si║[001]Al; (110)Si║(1-10)Al. The present work gives more information about such aluminum precipitates based on electron microscopy. A356 based alloys were cast in a steel chill mould. Both as-cast and solution-treated samples were studied by (scanning) transmission electron microscopy (STEM). Thin foils were prepared by ion polishing with high purity Ar+ at 4 keV. A JEOL-2010F microscope equipped with an energy dispersive x-ray spectroscopic (EDXS) detector was used. Linescan analysis in STEM mode was performed with a probe of 0.7 nm. Aluminum precipitates with diameters in a range of 10-25 nm were observed in both as-cast and subsequently heat treated samples (Figure 1) and their morphology and size were found to be insensitive to heat treatments. Additional alloying elements did not affect the precipitate composition. The shape of precipitate rich areas varied from one eutectic Si particle to another, but most of them showed a sharp boundary which separated precipitate rich areas from precipitate free areas. Bright field STEM image revealed precipitates including dark “head”. HRTEM images (Figure 2) and linescan analysis of selected precipitate showed that such dark “head” to be composed of Si. Oxygen impurities are postulated to gather aluminum and stabilize aluminum precipitates [5]. However, the present EDS analyses did not show higher level of oxygen from precipitates as compared to the Si matrix. It is suggested, based on the present results, that the formation of aluminum precipitates is probably related to the observed dark “head” [7]. 1. 2. 3. 4. 5.
L. Lu, K. Nogita, S. D. McDonald, A. K. Dahle, JOM 56 (2004) 52. H. Y. Kim, T. Y. Park, S. W. Han, H. M. Lee, J. Cryst. Growth 291 (2006) 207. Q. G. Wang, C. J. Davidson, J. Mater. Sci. 36 (2001) 739. S. C. Hogg, C. J. D. Hetherington, H. V. Atkinson, Phil. Mag. Lett. 80 (2000) 477. D. K. Sadana, M. H. Norcott, R. G. Wilson, U. Dahmen, Appl. Phys. Lett. 49 (1986) 1169.
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6. 7.
W. T. Donlon, Metall. Mater. Trans. 34A (2003) 523. The work has been conducted within the FONDAL project which is a cooperation NTNU, CNRS, Hydro Aluminium and Montupet SA. The authors would like to acknowledge financial support from the French-Norwegian Foundation for Scientific and Technical and Industrial Research as well as from Hydro Aluminium and Montupet SA.
Figure 1. (a) bright field TEM image showing a eutectic Si particle in A356 based alloy. Small precipitates are observed in a part region of Si particle. There is an obvious boundary separating precipitate area from unprecipitate area (marked by arrow). (b) intermediate magnification image from an area of (a) showing highly dense precipitates of 10-25 nm in diameter. (c) Bright field STEM image showing spheroidal morphology of precipitate including a dark “head”.
Figure 2. High resolution TEM images of precipitates taken along different zone diffractions (a) Si [111] and (b) Si [-4-8-9]. Moiré fringes are caused by different lattice constants of Al and Si structure. The dark “head” shows the same lattice fringe as the Si matrix. There is a detectable interface between Al-precipitate and its dark “head”.
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Microstructure of slow-cooled wedge-cast Cu58Co42 alloy with a metastable liquid miscibility gap E. Johnson1,2, S. Curiotto2, N. Pryds3 and L. Battezzati4 1. Niels Bohr Institute, Nano Science Center, University of Copenhagen, Denmark 2. Materials Research Department, RISØ DTU, Roskilde, Denmark 3. Fuel Cell and Solid State Chemistry Department, RISØ DTU, Roskilde, Denmark 4. Dipartimento di Chimica, IFM, Centro di Eccellenza NIS, University of Torino, Italy Erik Johnson, johnsonfys.ku.dk Keywords: Cu-Co alloys, Liquid immiscibility gap, Microstructure
The Cu-Co equilibrium phase diagram (figure 1) contains a metastable liquid phase miscibility gap where the liquid phase will separate into two, provided it can be undercooled sufficiently before solidification sets in [1]. This has been achieved by casting a Cu58Co42 alloy in a wedge-shaped mould where the as-cast microstructure showed evidence for liquid phase separation prior to solidification, e.g. formation of Co-rich droplets in a Cu- rich matrix [2,3]. Using backscattering SEM the dark Co-rich phase can easily be distinguished from the bright Cu-rich matrix (figure 2). The big dark particles are the Co-rich droplets that solidified after phase separation in the undercooled liquid. As cooling progressed in the liquid droplets the composition of the Co-rich phase would change and small droplets of Cu-rich liquid would form and be pushed towards the interior leaving the rim unaffected. TEM samples of the wedge-cast material proved to be difficult to produce by electropolishing due to the coarseness of the microstructure, and ion milling was found to be the only preparation technique that gave reasonable result albeit both phases were saturated with radiation damage in form of dislocation loops and tangles. Figure 3 shows low magnification bright-field TEM images from two different samples. The shape of the Co-rich particles is mostly spherical in accordance with their initial formation as liquid droplets. The larger Co-rich particles (figure 3a) can be recognized from the radiation damage structure that is denser than in the Cu-rich matrix which often contains a high density of semi coherent Co precipitates formed in secondary solid state reactions. The size of the smallest Co-rich particles appears to be as small as 50100 nm (figure 3b) while it is difficult to ascertain smaller particles due to the radiation damage induced dislocation loops. X-ray diffraction showed that both phases have fcc structure and that the lattice parameters deviated with less than 1%. This means that the two phases – Co-rich droplets and Cu-rich matrix - often are aligned and essential grow as larger single crystals. The interface boundaries are hence pure chemical interfaces with structural integrity across the boundaries. The small mismatch between the two phases is reflected in the misfit dislocation distributions in the interfaces (figure 3) that can easily be confused with moiré fringes (figure 2a).
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Y. Nakagawa: Acta Metallurgica, 6 (1958), p. 704. S. Curiotti, N. Pryds, E. Johnson, L. Battezzati, Mater. Sci. Eng. A, 449-451 (2007), p. 644. S. Curiotto, N. Pryds, E. Johnson, L. Battezzati, Met. Mater. Trans. A, 37A (2006), p. 2361. The work has been supported by the European Space Agency within the project “CoolCop” (ESA-MAP AO 99-010) and by the Danish Natural Science Research Councel.
1800 1700
s ta b le p h a s e d ia g ra m m is c ib ility g a p
γ
L iq u id
T [K]
1 6 0 0 C o -ric h 1500 1400 1300 1200 0 ,0
γ
C u -ric h
0 ,2
0 ,4 0 ,6 x (C u )
0 ,8
1 ,0
Figure 1. Cu-Co equilibrium phase diagram showing the occurrence of the metastable liquid miscibility gap. The arrow indicates the alloy composition.
Figure 2. SEM backscattering image from slow-cooled wedge-cast Cu58Co42 sample showing liquid phase separation that has led to formation of Co-rich droplets and particles dispersed in the Cu-rich matrix.
Cu-rich
1 μm
500 nm
Figures 3a and b. TEM of Co-rich particles formed by phase separation in the undercooled liquid in a Cu58Co42 slow-cooled wedge-cast sample. Most Co-rich particles have spherical shape.
100 nm
Co-rich
Figure 4. Interphase between a Co-rich particle and the Cu- rich matrix. The boundary separates phases with different composition but similar crystallography.
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TEM investigations of Elektron 21 magnesium alloy after long-term annealing A. Kielbus Silesian University of Technology, Katowice, Poland [email protected] Keywords: Elektron 21, magnesium alloy, long-term annealing, TEM, microstructure
The paper presents results of TEM investigations of Elektron 21 magnesium alloy in as cast condition and after long-term annealing. Elektron 21 is magnesium based casting alloy containing rare earth (Nd and Gd) for used to at 200°C in aerospace application. It has high strength, good corrosion resistance and excellent castability [1]. Magnesium alloys containing neodymium and gadolinium are characterised by high-strength and good creep-resistant alloys for automotive and aerospace applications [2,3]. The rare earth elements have beneficial effect of on the creep properties and thermal stability of structure and mechanical properties of magnesium alloys [4]. The strength of magnesium alloys with RE is achieved essentially via precipitation strengthening. These alloys precipitate from the solid solution according to the sequence of phases: αMg→β”→β’ →β [5]. The material for the research was a casting Elektron 21 magnesium alloy. The chemical composition of this alloy is provided in Table 1. Solution heat treatment was performed at 520°C/8 h/water. Ageing treatment was performed at 200°C/16h/air. Long-term annealing was performed at 350°C/100÷2000h and then quenched in air. The microstructure of Elektron 21 alloy was examined by a Philips CM 20 Transmission Electron Microscope. The microstructure of Elektron 21 consists of primary solid solution α grains with eutectic α + Mg12(Ndx,Gd1-x) phase at grain boundaries (Fig.1). The Mg12(Ndx,Gd1-x) phase is a modification of Mg12Nd phase with neodymium substituted by gadolinium due to reasonably small difference in the atomic radii. After solution treatment the Mg12(Ndx,Gd1-x) phase dissolved in the matrix. Aging at 200°C/16h leads to precipitation of β’ phase (Fig.2). After annealing at 350°C/100h the microstructure consists of β phase precipitates inside solid solution grains (Fig.3) and singular precipitates of Mg41Nd5 phase on grain boundaries. With continued ageing at 350°C (5000h) the volume fraction of Mg41Nd5 increase. The precipitates of this phase create the characteristic network on grain boundaries. Also the regular precipitates of Mg3Gd phase have been observed (Fig.4) 1. 2. 3. 4. 5.
P. Lyon, T. Wilks, I. Syed, Magnesium Technology (2005), p.303. H. Friedrich, S. Schumann, Journal of Materials Processing Technology, Vol. 117, Issue 3 (2001), p. 276. B. Mordike, Materials Science and Engineering, A324 (2002), p.103. B. Mordike, Journal of Materials Processing Technology, Vol. 117, Issue 3 (2001), p.391. B. Smola, I. Stulikova, F. von Buch, B. Mordike, Materials Science and Engineering A, 324, Issue 1-2 (2002), p.113.
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6.
The present work was supported by the Polish Ministry of Science and Higher Education under the research project No N N507 451334 “Structural stability of Mg-RE-Zr magnesium alloys in simulated working conditions”.
Table I. Chemical composition of the Elektron 21 magnesium alloy in wt.-%. Gd Nd Zr Zn Mn Fe Ag Mg 1,2 2,7 0,49 0,4 0,001 0,003 0,01 balance
Figure 1. Eutectic α + Mg12(Ndx,Gd1-x) in Elektron 21 alloy in as-cast condition.
Figure 2. β’ precipitates in Elektron 21 alloy after aging at 200°C/16h.
Figure 3. β precipitates in Elektron 21 alloy after annealing at 350°C/100h.
Figure 4. Mg41Nd5 and Mg3Gd precipitates in Elektron 21 alloy after annealing at 350°C/2000h.
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Microstructure of AJ62 magnesium alloy after long-term annealing A. Kielbus1 and J. Mizera2 1. Silesian University of Technology, Katowice, Poland 2. Warsaw University of Technology, Warsaw, Poland [email protected] Keywords: AJ62, magnesium alloy, long-term annealing, microstructure
The paper presents results of SEM investigations of AJ62 magnesium alloy in as cast condition and after long-term annealing. The AJ62 contains approximately 6% aluminum and about 2,5% strontium. The Mg-Al-Sr alloys exhibit good elevatedtemperature tensile properties, excellent creep resistance (better than Mg-Al-RE) and good castability [1]. The alloys have also Zn, Si And Mn. Strontium is added for creep resistance in die casting and gravity casting alloys or to reduced shrinkage porosity. Sr level may range from 0,02% to 3% [2]. The Mg-Al-Sr alloys (AJ52, AJ62) show different microstructure based on the Sr/Al ratio. For Sr/Al ratio below 0.3, the microstructure consists of α-Mg with precipitates of Al4Sr phase. When the Sr/Al ratio increases, a second intermetallic phase (probably Al3Mg13Sr) is observed together with Al4Sr phase. This ratio also controls the precipitation of Mg17Al12 phase. When Sr/Al ratio is very low this phase occur in the α-Mg matrix [3,4]. The material for the research was a sand casting AJ62 magnesium alloy. The chemical composition of this alloy is provided in Table 1. Sand casting was performed at 700°C temperatures. Annealing treatments were performed at 180°C, 250°C and 350°C during 500÷3000h with cooling in air. The microstructure of AJ62 alloy was examined by a Hitachi S3400 Scanning Electron Microscope. The AJ62 magnesium alloy after sand casting was characterized by the solid solution α with the lamellar eutectic (Al,Mg)4Sr + solid solution α and the globular precipitates of the Mn5Al8 phase. Morever the occurrence of not mumerous precipitates of the massive Al3Mg13Sr phase have been provided (Fig.1). Annealing this alloy at 180°C results the precipitation of Mg17Al12 phase and beginning of Al3Mg13Sr phase decomposition (Fig.2). Long-term annealing (3000h) at 250°C causes more far decomposition of Al3Mg13Sr phase (Fig.3). The precipitates of Mg17Al12 phase were not observed. Morphology and volume fraction of (Al,Mg)4Sr phase didn’t change. After exposure at 350°C the total decomposition of Al3Mg13Sr phase on mixture Al4Sr + αMg phases was observed (Fig.4). 1. 2. 3.
A.A.Luo, Materials Science Forum, Vols. 419-422, (2003), p. 57. H.E.Friedrich, B.L.Mordike, Magnesium Technology, Springer-Verlag Berlin Heidelberg, (2006). M.O.Pekguleryuz, Proc. 6th International Conference Magnesium Alloys and their Applications, Wolfsburg (2003), p.74.
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4. 5.
B.Jing, S.Yangshan, X.Shan, X.Feng, Z.Tianbai, Materials Science and Engineering A419 (2006), p.181. The present work was supported by the Polish Ministry of Science and Higher Education under the research project No PBZ-KBN-114/T08/2004.
Table I. Chemical composition of the AJ62 magnesium alloy in wt.-%. Alloy
Mg
Al
Mn
Sr
AJ62
balance
6,1
0,34
2,1
Figure 1. Precipitates of (Al,Mg)4Sr and Figure 2. Beginning of Al3Mg13Sr phase Al3Mg13Sr phases in AJ62 alloy after sand decomposition and precipitates of Mg17Al12 casting. phase in AJ62 alloy after annealing at 180°C/1000h.
Figure 3. Decomposition of Al3Mg13Sr phase after 3000h exposure at 250°C.
Figure 4. Total decomposition of Al3Mg13Sr phase on mixture Al4Sr + α phases after 1000h exposure at 350°C.
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EELS characterisation of the interface between nanoscaled ODS particles and matrix in advanced fusion steels M. Klimenkov, R. Lindau and A. Möslang Institute for Materials Research I, Forschungszentrum Karlsruhe GmbH Hermann-von-Helmholtz-Platz 1, 76344 Eggenstein-Leopoldshafen, Germany [email protected] Keywords: EELS, ODS steels, interfaces
Reduced activation ferritic–martensitic alloys (RAFM) are primary candidates for structural application in breeding blankets of fusion reactor plants, mainly due to their excellent swelling resistance and low activation properties. By the selection a of specific production route for ODS-Eurofer steel of the second generation, which included rolling and appropriate thermal treatments, DBTT could be shifted from values between +60 and +100°C for hipped ODS-Eurofer to values between -40 and -80°C° [1]. In addition ductility was improved significantly. This ductility improvement of the steel properties could be attributed to dislocation pinning on ODS particles. TEM examination of the ODS particle/matrix interface is a step toward understanding of the pinning mechanism and herewith the possibility to predict the steel properties. The TEM investigation of the ODS particle – matrix interfaces shows that ODS particles are in no case single Y2O3 phase dispersoids as it was suggested in the earlier publications [2]. The analytical TEM investigations using 1nm spatial resolution clearly show that all ODS particles have a complex shell/core structure (Figure 1). The Fe, Cr and Y - maps show, that the investigated Y2O3 ODS particle is imbedded in the matrix. The V-rich shell is clearly visible in the V-map. Performing spatially resolved EELS experiments clearly confirms the formation of a V rich shell around ODS particles (Figure 2). The intensity of the V-L2,3 edge increases at the edge of the ODS particle indicating a V-rich shell. The O-K edge shows the fine structure that is typical for Y2O3 material. The intensity of this edge is not zero in the regions outside the ODS particle due to the oxidation of specimen surface. Using the cross section scattering ratio of 1.8 for Mn-L2,3 and O-K EELS edges it was estimated, that the Mn/O ratio is about 0.05 to 0.10. From these considerations a (Y1.8Mn0.2)O3 composition of this ODS particle can be estimated. The earlier analytical investigations, which have been performed using EDX method did not detect Mn inside ODS particles because Mn-Kα and Mn-Kβ lines overlap with Cr-Kβ and Fe-Kα lines – the main alloying elements. The formation of thin oxide films on the specimen surface produces the weak O signal and disturbs herewith the typical energy loss near edge structure (ELNES) of measured spectra. This is especially important for the spectra obtained directly on the interfaces, where the ELNES fluctuations are very weak. From the spectra presented in the Figure 2 a “background” spectrum from the matrix was subtracted. The spectrum taken 1.5 nm from the interface in the matrix does not show any structure (Figure 3a)
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 443–444, DOI: 10.1007/978-3-540-85226-1_222, © Springer-Verlag Berlin Heidelberg 2008
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while the spectrum obtained directly on the interface (Figure 3b) shows V-L2,3 and O-K signals only. The signal from the Y2O3 core shows a strong O-K signal and Mn-L2,3 edge (Figure 3c). These difference spectra clearly show that the shell consists of V oxide. 1. 2.
R. Lindau, A. Moslang, M. Schirra, P. Schlossmacher, M. Klimenkov, Journal of Nuclear Materials 307-311 (2002) 769 M. Klimiankou, R. Lindau, A. Moslang Journal of Crystal Growth 249 (2003) 381
b)
c)
Fe
Cr
d)
e)
Y
V
a)
25nm
O-K
V-L2,3
Cr-L2,3
Mn-L2,3
1 .5
inte rfa c + e
+3.0
/nm
a)
-4.5
.5
Dis tan -3.0 ce fr -1 o m
0
b)
20 nm
500
30
600
Mn-L2,3
Figure 3. The interface spectra after subtraction of background spectrum from the surrounded matrix.
O-K
counts
Cr-L2,3 c)
20 10
V-L2,3 O-K
b) a)
0
550
600
Energy /keV
Figure 2. Spatially resolved EELS investigations of ODS particle /matrix interfaces. The area with two ODS particles is imaged in the part (a). The row of EELS spectra obtained while the scanning along the marked line is presented in part (b).
550
Energy /keV
650
650
Figure 1. EDX elemental mapping of an area with ODS particle: (a) shows the HAADF image, (b) – (e) represent the elemental maps obtained using Fe-Kα, Cr-Kα, Y-L, and V-Kα EDX lines respectively.
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Microstructure-mechanical property relationships in a maraging 250 steel P. Kokkonidis, E. Papadopoulou, A. Rizos, T. Koutsoukis and G. Fourlaris Laboratory of Physical Metallurgy, National Technical University of Athens, 9 Heroon Polytechniou St. , 15780, Athens Greece [email protected] Keywords: maraging 250, microstructure, mechanical properties, ageing
The objective of the present study is to examine the relationship of the microstructure with its mechanical properties in a maraging 250 steel, following a series of selected ageing experiments. Furthermore, the effect of cold plastic deformation (cold rolling, at different amounts of cold rolling reduction) on the mechanical properties of the steel following ageing was also studied. Maraging steels represent a distinct family of high strength steels. They exhibit a characteristic strengthening sequence following a precipitation at 480 o C [1]. A typical characteristic of this family of steels is their combination of enhanced ductility with high tensile strength values. [2] The samples were initially solution treated at 820°C for 30mins followed by water quenching (room temperature). Aging took place at 460°C, or 480°C or 510°C for periods of time ranging from 30mins to 24h, followed by water quenching (20οC). For the microstructure investigation a scanning electron microscope (Philips XL30) was employed coupled with X-ray diffraction phase identification. The hardness of specimens was measured using a Vickers hardness testing facility. Finally, selective tensile testing was carried out in specimens that were aged at 510οC. The experimental results showed an increase in hardness at all three chosen ageing temperatures, even for the brief ageing time of 30 mins. Aged specimens were compared with specimens that were annealed at 820ο C for 30 min and quenched in water. Following prolonged ageing, a decrease in hardness was observed for ageing times of more than 24h at 510οC (Figure 1). Tensile testing confirmed the influence of ageing on the mechanical properties of this steel grade (Figure 2). Increased hardness occurs due to the precipitation reactions of secondary phases that take place during the ageing process. Precipitation of reverted austenite combined with the coarsening of primarily fine precipitates causes a marked decrease in hardness at an ageing temperature of 510οC. Increased values of hardness were observed on specimens subjected to cold plastic deformation prior to ageing at 485οC at brief aging intervals, however, these results were influenced by the amount of cold deformation applied. Typical aged microstructures are given in the SEM micrographs of Figure 3. The evolution of various phases identified by XRD is presented in Figure 4. 1. 2.
Michael Schmidt and Kurt Rohrbach, “Heat Treating of Maraging Steels”, ASM Metals Handbook, Volume 4, (1991) U. K. Viswanathan, G. K. Dey, V. Sethumadhavan: “Effect of Austenite Reversion during Overageing on the Mechanical Properties of 18 Ni (350) Maraging Steel”, Materials Science and Engineering A, vol 398 (2005) p. 367-372
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Figure 1. Comparative evolution of hardness Figure 2. Mechanical properties of the during ageing versus ageing time ageing maraging steel versus ageing time at a temperature of 510 oC
Figure 3. Secondary electron micrograph of a) Solution treated at 820 oC b) aged at 510 o C for 24h.
Figure 4. XRD traces at an ageing temperature of 510 oC for various aging times
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Microstructure of Co-Ni based superalloys T.J. Konno, T. Tadano, H. Matsumoto and A. Chiba Institute for Materials Research, Tohoku University 2-1-1 Katahira, Aoba-ku, Sendai, 980-8577, Japan [email protected] Keywords: microstructure, Co-Ni alloy, HRTEM
Co-based alloys are used in numerous commercial applications because of their excellent performance under severe environments, which require high-temperature strength, corrosion resistance, wear resistance, etc. In general, these alloys belong to a multi-component system, and the composition and heat and/or mechanical treatments to achieve a desired property are often complex. In this respect, understanding the microstructural changes during the processing is indispensable for the optimal selection of various processing parameters. In particular, the identification of phases in the early stages of phase transformation, which necessitates atomic level observations, is critical to elucidate the origin of the mechanical properties of the alloys. In the present study, we report the precipitation behavior of a Co-Ni-Cr-Mo alloy, heat-treated at 973-1073K, which results in a mechanical strength suitable for hightemperature applications. For transmission electron microscopy (TEM) work, we used Titan 80-300 operating at 300kV, equipped with a Cs corrector for the objective lens; and employed a range of techniques available to clarify the crystallography and kinetics of the phase transformation. They include STEM-BF/HAADF imaging, EDS mapping and 3D tomography, in addition to conventional high-resolution microscopy (HRTEM). Figure 1 is a STEM-HAADF micrograph of the alloy annealed at 1073K for 3 weeks, showing well lamellar precipitates. Simultaneous EDS analysis suggested that Ni and Mo are enriched in the precipitates, whereas Cr is partitioned out from them. Diffraction patterns from these precipitates suggested they belong to a hexagonal system with the basal plane parallel to the {111} planes of the matrix. 3D tomography also supported this view. Figure 2 is a HRTEM micrograph of the interface region, viewed from the [110] direction of the matrix fcc phase, showing that these precipitates possess a well-defined coherent interface with the matrix, although conventional TEM showed contrasts in the surrounding matrix region due to strain fields. The c-axis of the hexagonal precipitate is close to 2/3 amatrix as expected from the simple ABAB... stacking sequence, whereas the a-axis is twice as much as the value expected from the sequence. This was also confirmed from the corresponding diffraction pattern and STEM images. This result shows that the precipitates are formed by an ordered arrangement of Mo and/or Ni atoms. 1. 2.
A. Chiba, X.G. Li and M.S. Kim Philos. Mag. A79 (1999) p. 1533-1554. A. Chiba and M.S. Kim Mater. Trans. 42 (2001) p. 2112-2117.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 447–448, DOI: 10.1007/978-3-540-85226-1_224, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. STEM-HAADF image of Co-Ni-Cr-Mo alloy annealed at 1073K for 3 weeks, viewed from [110] zone axis of the fcc matrix.
Figure 2. High-resolution micrograph of an interface region of the f.c.c. matrix and a hexagonal precipitate, viewed along [110] zone axis of the matrix. Note the double periodicity in the direction parallel to the interface, i.e., the [210] direction of the hexagonal precipitate.
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Precipitation reactions in superferritic stainless steels T. Koutsoukis, K. Konstantinidis, P. Kokkonidis, E. Papadopoulou and G. Fourlaris Laboratory of Physical Metallurgy, National Technical University of Athens, 9 Iroon Polytechniou Street, 15780, Athens, Greece [email protected] Keywords: superferritic stainless steel, intermetallic, precipitation.
Superferritic stainless steels exhibit an excellent combination of corrosion resistance, especially under chloride stress corrosion cracking conditions, and desirable mechanical properties at elevated temperatures. The lower amounts of Ni makes these grades a cheaper solution in relation to superaustenitic grades, while maintaining strength characteristics with modest toughness. Hence, they are widely used in heat exchanger tubing and thin sheet applications. Precipitation or transformation of secondary phases takes place when superferritic stainless grades are exposed at elevated temperatures, adversely affecting their mechanical properties and corrosion behaviour [1,2,3]. The purpose of the present paper is to study the phase transformations occurring during ageing of superferritic stainless steels. Thus, a novel superferritic grade (SFSS), produced via the HIP method with controlled alumina additions, was employed and subjected to 10% and 20% cold rolling. Composition of the steel studied is shown in Table I. Samples were heat treated within the temperature range of 650oC to 950oC, for times up to 500h, followed by water quenching. The evolution of microstructure was assessed using a Philips XL30 SEM microscope, while EDS microanalysis and XRD diffraction were utilized to characterise the nature and chemical composition of the microstructure components observed. Finally, Vickers hardness testing was performed on all samples. Several intermetallic phases were identified following ageing of superferritic material within the temperature range of 650oC to 950oC. Sigma (σ) phase was the most common secondary phase observed (Figures 1A, 1B). Following ageing for 24h at 650oC sigma (σ) phase was observed forming primarily at grain boundaries. The volume fraction of sigma (σ) phase increases in proportion to ageing time and temperature. The same transformation behaviour was observed following ageing at 750oC and 850oC but exhibiting faster transformation kinetics. Sigma (σ) phase formation was noticed following 24h or 1h of ageing at 750oC or 850oC, respectively. On the contrary, slower transformation kinetics were observed following ageing at 950oC. Sigma (σ) phase was identified by XRD analysis and EDS spot microanalysis performed on sigma phase particles indicating a high Cr content. In addition to sigma (σ) phase, a second precipitate species, rich in Mo, was observed at grain boundaries of both cold deformed samples (Figure 1A). This phase is likely to be chi (χ) phase [2,3] or one of the group of Laves phases [1]. Finally, in figure 2 the evolution of hardness (HV) versus ageing time (h) for both 10% and 20% cold rolled superferritic steels is shown, for various ageing temperatures.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 449–450, DOI: 10.1007/978-3-540-85226-1_225, © Springer-Verlag Berlin Heidelberg 2008
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1. 2. 3.
De Andrade TF et al, Article in Press, d.o.i. 10.1016/j.matchar.2007.03.006. P. G. Ng, E. Clarke, C. A. Khoo and G. Fourlaris, Mat. Sc. and Tech., 22 (2006), p. 852. D. M. E. Villanueva, F. C. P. Junior, R. L. Plaut and A. F. Padilha, Mat. Sc. and Tech., 22 (2006), p. 1098.
Table I. Nominal composition of steel studied (wt %). C 0.010
SFSS
A σ
Si 0.27
Mn 1.11
Cr 25.17
Ni 2.63
Mo 3.84
N 0.017
Cu 0.10
B
α
Al 0.01
α
Mo-rich
σ
Figure 1. A. Secondary electron image of 10% cold rolled superferritic stainless steel aged at 650oC for 500h showing sigma (σ) phase formation and Mo-rich (possibly chi) phase precipitation and B. Backscattered electron image of 20% cold rolled superferritic stainless steel aged at 850oC for 500h showing sigma (σ) phase formation. 600
o
20%/750 C
500
o
o
Hardness (HV)
20%/850 C
400
10%/750 C
o
10%/850 C o
10%/950 C
300
o
10%/650 C o
20%/650 C
200
o
20%/650 C
100 0
100
200 300 Ageing Time (h)
400
500
Figure 2. Evolution of hardness (HV) versus ageing time (h) for 10% and 20% cold rolled superferritic stainless steels, for various ageing temperatures.
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Effect of ageing in cold rolled superaustenitic stainless steels S. Zormalia, T. Koutsoukis, E. Papadopoulou, P. Kokkonidis and G. Fourlaris Laboratory of Physical Metallurgy, National Technical University of Athens, 9 Iroon Polytechniou Street, 15780, Athens, Greece [email protected] Keywords: superaustenitic stainless steel, deformation, precipitation.
Superaustenitic stainless steels exhibit excellent corrosion resistance, at room, sub zero and elevated temperatures, especially in chloride containing environments, coupled with excellent toughness and weldability. These steels are prone to precipitation of secondary phases when exposed at elevated temperatures, affecting their mechanical properties and corrosion behaviour [1,2]. The purpose of the present paper is to study the combined effects of cold deformation and heat treatment in the microstructure and mechanical properties of superaustenitic stainless steels. Two superaustenitic grades, S32654 and S31254, were used for this study. The nominal composition of the steels studied is shown in Table I. Samples of both steels were cold rolled to 20%, 40% and 60% reduction in thickness and then heat treated within the temperature range of 650oC to 950oC, for various times up to 120h, followed by water quenching. The evolution of microstructure was assessed using optical microscopy and SEM (Jeol 6380LV) microscopy coupled with EDS microanalysis. Finally X-ray diffraction was used to characterise the evolution of phase transformation. Vickers hardness testing was also performed on all samples. Several precipitation species were identified following ageing of superaustenitic steels studied within the temperature range of 650oC to 950oC. Sigma (σ) phase was the dominant intermetallic phase observed (Figures 1A, 1B). Micro-segregation was observed in all cold rolled samples of S31254 before ageing (as reference condition) mainly at the centre of the samples and parallel to rolling direction. Sigma (σ) phase precipitates were identified following ageing for 1h at 750oC, 850oC and 950oC at all deformation studied on both steel grades. Following ageing at 650oC precipitation of sigma (σ) phase begins approximately at 24h. Precipitates form primarily at grain boundaries and triple point junctions but also longitudinal on prior slip planes visible following the cold rolling. Volume fraction of sigma (σ) phase increases in proportion to the amount of deformation, ageing time and temperature. Sigma (σ) phase formation was identified by XRD phase analysis. In addition to sigma (σ) phase, a second precipitate species, rich in Mo, was observed at grain boundaries on all deformed states (Figure 1A). This species following XRD analysis is hypothesized to be either Laves phase (Laves-η phase) or chi (χ) phase [1,2]. Figure 2 shows the typical evolution of hardness (HV) versus ageing time (h) for 60% cold rolled S32654, for various ageing temperatures. At the present study many factors affect the hardness of the steels studied. The amount, shape and size of all
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 451–452, DOI: 10.1007/978-3-540-85226-1_226, © Springer-Verlag Berlin Heidelberg 2008
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precipitation species is an important factor that in general results to higher hardness values (precipitation hardening) which are enhanced by prior cold deformation. Ageing following cold deformation can enhances the magnitude of burl precipitation reaction and minimizes time required for initiation of precipitation reaction. 1. 2.
J. Vrest’al, A. Kroupa, M. Sob, Com. Mat. Sci., 38(2006), 298-302 R. W. Fonda et al., Met. Mat. Trans. A, 38A(2007), 2721.
Table I. Nominal composition of steels studied (wt %). C
Si
Mn
P
S
Cr
Ni
Mo
N
Cu
Ti
S32654
0.013
0.24
3.43
0.021
0.001
24.19
21.58
7.24
0.497
0.38
0.001
S31254
0.012
0.36
0.47
0.019
0.001
20.02
18.16
5.98
0.214
0.65
0.001
A
σ
Mo rich
B
σ
Figure 1. A. SEI micrograph of 20% cold rolled S32654, aged at 850oC for 120h showing sigma (σ) phase precipitation and B. Backscattered electron image of 60% cold rolled S31254 aged at 850oC for 120h showing sigma (σ) phase (black contrasted) and Laves (η) phase (white contrasted) precipitation. 550
AR
Hardness (HV20)
500
o
650 C
450
o
750 C
400
o
850 C
o
950 C
350 300 250 200 0
12
24
36
48 60 72 Ageing Time (h)
84
96
108
120
Figure 2. Evolution of hardness (HV) versus ageing time (h) for 60% cold rolled S32654, for various ageing temperatures.
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Structure and properties of P/M material of AlMg – SiO2 system processed by mechanical alloying A. Kula1, L. Błaż1, M. Sugamata2, J. Kaneko2, Ł. Górka1, J. Sobota1 and G. Włoch1 1. University of Science and Technology, av. Mickiewicza 30, 30 – 059 Cracow, Poland 2. Nihon University, Chiba 275 – 8575 Japan [email protected] Keywords: Mechanical alloying, silicon oxide, Aluminum matrix composite, Transmission Electron Microscopy (TEM), Scanning Transmission Electron Microscopy (STEM)
New methods of material processing have been actively pursued in recent years in an attempt to extend current materials performance. Of particular interest are the powder metallurgy (P/M) techniques of mechanical alloying (MA). The MA process is generally used to create the materials with unique properties which give the material a wide spectrum of possible advanced applications. Light - metal based mechanically alloyed composites strengthened by heavy metal oxides addition (MeO) [1] have been examined according to bilateral research cooperation between Nihon University, Tokyo and AGH – University of Science and Technology. In this paper, a powder of Al – 10% at Mg alloy was mechanically alloyed with addition of silicon oxide (SiO2). Metal oxide was added by such an amount that the constituent oxygen corresponds to 6 at % of a total charge of AlMg – oxide mixture. MA processing of powdered components was carried out by using an Attritor ball milling for 30 hours with addition of 5 % methanol that was added to the powders as the process control agent. MA powders were then consolidated by cold pressing, vacuum degassing and hot extrusion at reduction ratio 25 to 1. Applications of that consolidation route ensure low porosity of extruded material. The composite was manufactured at Nihon University laboratories in Tokyo. Because of fine microstructure attained by MA, solid state reactions at high enough temperature are promoted due to increased reaction area, decreased diffusion distance and enhanced diffusivity of elements. As a result, Mg2Si, MgO and Al/Mg-oxides are formed in result of SiO2 decomposition in the AlMg matrix. It is worth stressing that decomposition of SiO2 was not observed in Al – SiO2 system [2]. It is clear that AlMg alloy can provide higher force for reduction of the SiO2 than pure Al. Typical microstructure of as – extruded and long - annealed sample are shown in Fig. 1a , b respectively. Long annealing time and favorable diffusion conditions were found to result in significant coarsening of Mg2Si phase (Fig 1b). Reduction of AlMg – SiO2 composite microhardness observed for samples annealed at 773K was negligible. It was ascribed mainly to the effect of disperse nano-sized Al(Mg)-oxides that retarded the coarsening of very fine structure of the material during prolonged annealing. In order to test the temperature effect on the composite mechanical properties, hot compression test were performed in the temperature range of 293 – 823 K. Typical set
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 453–454, DOI: 10.1007/978-3-540-85226-1_227, © Springer-Verlag Berlin Heidelberg 2008
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of stress – strain curves is shown in Fig. 2. High flow stress values were ascribed to highly refined structure and low porosity of the material that was remained in spite of chemical reaction between composite components. 1. 2.
L. Blaz, J. Kaneko and M. Sugamata, Material Chemistry and Physics 81 (2003), 387 - 389 D.G. Kim, J. Kaneko and M. Sugamata, J. Japan Inst. Metals, 57 (1993), 1325
IP
1 μm
2 μm a)
b)
Figure 1. STEM micrographs of AlMg – SiO2 composite; a) as – extruded material, b) annaealed at 773 K for 168h; IP – intermetalic phase Mg2Si
a)
b)
Figure 2. (a) Stress – strain curves for AlMg – SiO2 composite. Temperature of deformation is marked on the figure. Hot compression test was performed at constant strain rate 5*10-3 s-1. (b) Effect of deformation temperature on the maximum flow sterss for as extruded samples. Some results for another Al – based composites are shown for comparsion purposes (to be published).
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Extrusion of rapidly solidified 6061 + 26 wt% Si alloy A. Kula1, M. Sugamata2, J. Kaneko2, L. Błaż1, G. Włoch1, J. Sobota1 and W. Bochniak1 1 University of Science and Technology, av. Mickiewicza 30, 30 – 059 Cracow, Poland 2 Nihon University, Chiba 275 – 8575 Japan [email protected] Keywords: Al – Si alloys, grain refining, Rapid Solidification (RS), Transmission Electron Microscopy (TEM), Scanning Transmission Electron Microscopy (STEM)
Due to low thermal expansion coefficient, high strength to weight ratio as well as high wear and corrosion resistance, high silicon aluminum alloys are commonly used in aerospace, automobile industry and many other practical applications [1]. The most common aluminum foundry alloys contain 5 – 12 wt % silicon. Higher content of silicon is not useful for commonly produced as-cast materials because of embrittlement effect resulted from the coarse primary silicon development. The most of mechanical properties of castings are determined by the silicon and eutectic structure morphology [2]. Some refining of the structure can be achieved by means of mechanical mold vibrations at high enough amplitudes [3], modifications with alkali or rare earth metals [4] and also by increasing the cooling rate during casting procedures [2]. However, the most effective refining of Al – Si alloy structure can be achieved due to rapid solidification (RS) combined with powder metallurgy (P/M) methods. Experiments described bellow were performed on RS powder of 6061 + 26 wt% Si alloy that was mechanically consolidated by vacuum hot compression and extrusion procedures. The RS powder was produced by air spray atomization at Toyo Aluminum Company. Chemical composition of the alloy is shown in Table 1. Table 1. Chemical composition of 6061+26%Si alloy Element Si Mg Cu Fe Cr wt.%
26.0
0.69
0.16
0.21
0.11
Zn
Al
0.11
in balance
The RS-powder was cold pressed in thin-wall 6061 alloy cans and degassed under pressure of ~10-4 Pa. As compressed billets were extruded by means of KOBO® method with cross-section reduction of λ=19 [5]. The method was also used for extrusion of tubes with cross-section reduction of λ=33. Typical structure of as – extruded material is shown in Fig. 1a. Bimodal distribution of quasi – spherical Si-particles was observed. Silicon particles were revealed by means of STEM element mapping and EDX analysis system (Fig.1b). It was found that prevailing silicon size was less than 0,7 µm in diameter. The maximum on the silicon histogram (Fig. 2) was observed within the range of 0,2 – 0,4 µm. Hot compression tests performed on as – extruded material within the temperature range of 293 – 823 K revealed low ductility and high strength of material. For comparison purposes the samples annealed at 773K/30 min were also tested by hot S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 455–456, DOI: 10.1007/978-3-540-85226-1_228, © Springer-Verlag Berlin Heidelberg 2008
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compression test. Fracture of hot deformed samples was not observed within used strain range, εt ≈0,4. The effect of deformation temperature on the flow stress maximum is shown in Fig. 3. The annealing of the samples resulted in the material recovery and following reduction of the flow stress value during compression test. Silicon morphology was not practically affected by preliminary heat treatment and hot deformation procedures. 1. 2. 3. 4. 5. 6.
K.U. Kainer, Metal matrix composites, Edit. K.U Kainer Weinheim, 2006 M.M Makhlouf and H. V. Guthy, J. Light Metals 1 (2001) 199 X. Jian, T.T. Meek and Q. Han, Scripta Materialia 54 (2006) 893 A. Knuutien, K. Nogita, S.D. McDonald and A.K. Dahle, J. Light Met. 1, (2001) 229 W. Bochniak and A. Korbel, Mater. Sci. Forum 331-3 Part 1-3, (2000), 613 S. Piotrowski, “Siluminy tlokowe” (in Polish), edit. PAN Katowice (1997)
a) b) Figure 1 (a) STEM picture of as – extruded material; (b) Bimodal Si-particle distribution revealed by STEM and EDX element mapping method
Figure 2. Silicon particles size distribution
Figure 3. Effect of deformation temperature on the maximum flow stress. Tension test result (1) –Y.S. and U.T.S. for conventional cast Al-25%Si-1%Cu1%Ni-1%Mg-0,4%Cu-0,5%Fe alloy are marked for comparsion purposes [6]
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TEM and EELS study of carbide precipitation in low alloyed steels C. Leguen1, M. Perez1, T. Epicier1, D. Acevedo2 and T. Sourmail2 1. Université de Lyon; INSA-Lyon, MATEIS, umr CNRS 5510, bât. B. Pascal, F-69621 Villeurbanne Cedex 2. ASCOMETAL CREAS-Metallurgy, F-57301 Hagondange Cedex [email protected] Keywords: precipitation, carbides, nitrides, steels, HRTEM, EELS
Microalloyed steels have received considerable interest over many years because of their extensive use for many industrial applications. As an example, the addition of vanadium and/or niobium is a well-known way to control the mechanical properties of such materials. In this context, understanding the evolution of the precipitation state during the elaboration process of steels is a key to optimising its final properties. From an experimental point of view, it is then required to proceed to a detailed microstructural characterization of the size, volume fraction, chemistry and crystallography of the precipitates. Generally speaking, observation of thin foils is needed to approach the real distribution, location, and orientation relationship of the precipitates within the matrix. However, extraction replicas remain the best way to determine the chemical composition. EELS (Electron Energy Loss Spectroscopy) is then required for accurate measurements of the carbon and nitrogen contents. For that purpose, aluminium or aluminium oxide replicas, free of any undesirable carbon, have been used. HRTEM and EELS have been applied to various low alloyed steels, e.g. a model Fe-V-C and more complex Fe-Nb-V-Ti-Al-C-N systems. Two illustrations are given in the following. Figure 1 concerns the crystallography of vanadium carbide precipitates. It has been observed that vanadium carbides nucleate from a super-saturated solid solution with an ordered V6C5 structure both in ferrite and in austenite. Indeed, the V6C5 superstructure has been positively identified when nucleation occurs in ferrite (Fig. 1a). Same observations have been obtained after a treatment of 1 hour at 1280°C followed by 10 days at 800°C, leading to nucleation in austenite. At least for the material investigated, this proves that vanadium carbides nucleate with the V6C5 ordered structure, contrary to previous literature results as discussed in [2]. The second illustration concerns EELS analyses performed on particles extracted from a Fe-Nb-V-Ti-Al-N-C alloy (figure 2). The fitting procedure used to determine the chemistry, and involving a combination of references spectra (Nb, V, Ti, C and N), will be detailed [3]. 1. 2. 3.
T. Gladman, Physical Metallurgy of Microalloyed Steels (The Inst. of Mater., London, 2002). T. Epicier, D. Acevedo, M. Perez, Philos. Mag., 88, 1, (2008), 31. The CLYM (Centre Lyonnais de Microscopie) is gratefully acknowledged for the access to the 2010F microscope.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 457–458, DOI: 10.1007/978-3-540-85226-1_229, © Springer-Verlag Berlin Heidelberg 2008
a)
458
b)
5 nm Figure 1. HRTEM images of vanadium carbide (observed along a ‹110›fcc zone axis) precipitated within FeCV alloy (after a treatment of 15 hours at 700°C). a): HRTEM micrograph showing a contrast re-enforcement every two (111)fcc lattice fringes due to carbon ordering in a V6C5 structure: the numerical diffractogram showing the superlattice reflection responsible for the fringe doubling (arrows). b): Same as a) after some seconds under the electron beam : note that the superlattice fringes and reflection have vanished due to the irradiation damage.
a)
c)
1 2
1
C-K N-K
25 nm
Ti-L2,3 Nb-M4,5
2
b) Figure 2. Complex carbide analyzed on an Al-extraction replica. a-b): conventional TEM and HAADF (High Angle Annular Dark Field) images; c): EELS spectra from the circled zones (top 1 and bottom 2 respectively) in a), showing that the upper part (arrow in b) of the crystal is TiN-rich, while the largest lower part is essentially NbC-based.
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Microstructural analysis of plastically deformed complex metallic alloy κ-AlMnNi M. Lipińska-Chwałek1,2, M. Heggen2, M. Feuerbacher2 and A. Czyrska-Filemonowicz1 1. Faculty of Metals Engineering and Industrial Computer Science, AGH University of Science and Technology, Al. Mickiewicza 30, 30-059 Kraków, Poland 2. Institut für Festkörperforschung, Forschungszentrum Jülich GmbH, D-52425 Juelich, Germany [email protected] Keywords: complex metallic alloys, dislocations, plastic deformation
Complex Metallic Alloys (CMAs) are an exceptional group of intermetallic phases possessing “giant” unit cells comprising some ten to some thousand atoms per unit cell [1]. One representative of this group of materials is the phase κ-AlMnNi, which was recently discovered in the Al-Mn-Ni system [2]. It is hexagonal (space group P63/m) with lattice parameters a = 1.76 nm and c = 1.24 nm, and has approximately 227 atoms in the unit cell. Recently we have successfully grown single-crystals of the κ-AlMnNi by means of self-flux growth technique. From these crystals, rectangular samples of about 1.2 x 1.2 x 3.0 mm3 were cut for plastic deformation experiments. The tests were performed in uniaxial compression along the [0001] direction at temperatures above 700°C with a constant strain rate of 10-5 s-1. Microstructural characterization of the deformed samples was performed by transmission electron microscopy (TEM). Dislocation motion on {1 1 00} planes and on (0001) basal planes was observed. The majority of dislocations move on {1 1 00} planes. These dislocations are loop-shaped partials elongated along the [0001] direction. Their Burgers vectors are found to be parallel to the < 1 1 00 > directions, i.e. perpendicular to the habit plane of the dislocations. The dislocations thus are prismatic edge loops moving by pure climb. Figure 1 shows such a dislocation loop imaged under bright-field two-beam conditions. Despite the fact that all used reciprocal vectors fulfil G G the extinction condition g ⋅ b = 0 , the loop segments often show a strong residual contrast. Full extinction can be achieved for those segments, for which the condition G G G g ⋅ b × l = 0 is fulfilled. A second set of prismatic edge loops was found on the (0001) basal planes. The dislocations were identified as partials with Burgers vectors parallel to the [0001] direction, i.e. their motion takes place by pure climb. Similarly to the dislocations in {1 1 00} habit planes, a complete extinction of the dislocation segments was observed if G G G the condition g ⋅ b × l = 0 is fulfilled. The present investigations show that pure climb of partial edge dislocations on basal (0001) and {1 1 00} planes is the primary deformation mechanism in the κ-AlMnNi phase.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 459–460, DOI: 10.1007/978-3-540-85226-1_230, © Springer-Verlag Berlin Heidelberg 2008
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K. Urban, M. Feuerbacher, J. Non-Cryst. Sol. 334-335 (2004) 143 S. Balanetskyy, G. Meisterernst, M. Feuerbacher, Intermetallics (2007), submitted The authors acknowledge a financial support from the European Network of Excellence (NoE) Complex Metallic Alloys, contract no NHP-CT-2005-500140.
a
c
b
d
200nm Figure 1: TEM bright-field micrographs of dislocations in the (1-100) habit plane in deformed κ-AlMnNi imaged using the reflections [ 44 8 0] (a), [000 4 ] (b), [ 4 480] (c) and [ 44 8 2] (d).
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Electron microscopy analysis of Mn partitioning in retained austenite- martensite- bainite islands A. Lis1, J. Lis1 and P. Wieczorek1 1. Lab of Electron Microscopy, Institute of Materials Engineering, Częstochowa University of Technology, 19 Armii Krajowej, 42-200 Częstochowa, Poland [email protected] Keywords: analysis, Mn partitioning, austenite- martensite- bainite islands
The partitioning of Mn in the multiphase microstructure of 6Mn16 steel after annealing at 973K for 1800 s was investigated using transmission electron microscopy (TEM; Philips CM20) with energy- dispersive X- ray spectrometry (EDS – EDAX). In previous paper it was established that increasing annealing time below AC1 increases the concentration of Mn in cementite and thus in new- formed austenite during next annealing in the (α+γ) temperature range [1]. Islands of the new- formed austenite at 973 or 1023 K will have different chemical composition depending on nucleation sites at interfaces of α/α or α/cementite. In the case of austenite nucleation at interface boundary α/α, initial Mn concentration C0Mn is higher than maximum solubility of Mn in ferrite at temperatures 973 or 1023 K - Cα0Mn. Initial Mn concentration C0Mn in austenite is much lower than maximum solubility of Mn at given temperature CγTMn. Diffusion rate of Mn in ferrite is one order faster than in austenite. DαMn = 7.12 x 10-13 cm2/s at 973 K. Thus Mn which diffuses away from ferrite will form enriched in manganese region at the α/γ interface. This was observed in the vicinity to austenite island which was transformed on cooling in the middle into bainitemartensite area presented in “Figure 1 a.”, while surroundings formed retained austenite. The increased concentration of Mn was confirmed by mapping analysis. During time of annealing at given temperature in (α+γ) range, Mn moves from ferrite to austenite across interface surface between austenite and ferrite until equilibrium state of the chemical potentials of Mn in austenite and ferrite will be reached. The balance depends on temperature and time of annealing. Small, stable, retained austenite islands were also detected. The example is shown in “Figure 1 b”. The high concentration of Mn in the austenite island is visible. When nucleus of γN is formed at the interface surface of the ferrite/alloyed cementite, then two interface surfaces are created i.e. austenite/carbide and ferrite/austenite. On each of their boundaries, different concentration of Mn will arise which are dependent on temperature, concentration of alloying elements in phases, degree of curvature of surfaces and volume ratio of carbides. Thus in multiphase microstructure there are different concentrations of Mn due to partitioning effects. In “Figure 1c.” the multiphase area is presented. Ferrite has 2.8 %Mn, bainitemartensite region 5.4 % Mn and bainitic ferrite 3.8 % Mn.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 461–462, DOI: 10.1007/978-3-540-85226-1_231, © Springer-Verlag Berlin Heidelberg 2008
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J. Lis, J. Morgiel and A. Lis, Materials Chemistry and Physics 81 (2003) pp.466-468.
a)
Mn
b)
Mn
c)
Ferrite – 2.8%Mn
Bainite Ferrite – 3.8%Mn
Bainite -martensite – 5.4%Mn
Figure 1. a) Bainite- martensite- retained austenite island b) Retained austenite island c) Multiphase microstructure with EDS spectra
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TEM characterization of microstructures in a Ni2MnGa alloy H. Maeda1, E. Taguchi2, K. Inoue3 and A. Sugiyama4 1. Fac. Sci. & Tech. Ryukoku Univ., 2. Research Center for UHVEM Osaka Univ., 3. Advanced Res. Inst. Waseda Univ., 4. Dept. Entrepreneur Osaka Sangyo Univ. [email protected] Keywords: Ni2MnGa, magnetostriction, off-stoichiometric alloy, TEM, microstructure
Since the appearance of huge magnetostriction in a Ni-Mn-Ga off-stoichiometric alloy was observed in 1996 [1], a number of researches have been made on this subject [2]. The off-stoichiometric alloys have a tetragonal based Heusler structure, and have both its Martensitic transformation temperature and its Curie temperature near the room temperature. In the bulk of Ni2.18Mn0.82Ga1.00, the reversible magnetostriction has been observed at room temperature [3] . In order to get information on microstructures in this alloy, TEM observation was carried out. The composition of TEM foils as determined by EDX was Ni2.22Mn0.80Ga0.98. A microstructure with stripes (domains) was observed. Figure 1 shows a selected area diffraction pattern obtained. Two sets of diffraction net patterns (i.e., T1 and T2) in a nearly twin relation were recognized in this figure. The incident beam direction was [010]T1 and [0 1 0]T 2 , and was vertical to the normal of the T1/T2 interfacial plane(i.e., (10 1 ) T1 and (10 1 ) T 2 ). From the distances between 000 and 400T1 and between 000 and 004T1, lattice parameters a and c for domain T1 can be estimated to be 0.556 and 0.664 nm, respectively. The c / a is then given as 1.194. These values are in good agreement with those obtained by the neutron diffraction [3]. It is noted here that a slight change in lattice parameters was observed between domains T1 and T2, under a constraint of constant a 2 + c 2 . Figure 2 a) and b) show dark field images obtained with a diffraction spot marked with a white and a gray circle in Figure 1 a), respectively. It is evident that the sample is composed of domains T1 and T2 which are alternatively stacked in the [10 1 ]T1 (= [10 1 ]T 2 ) direction. Figure 2c) is a schematic illustration showing the orientation relationship between the incident beam and the sample. 1. 2. 3. 4.
K.Ullakko, J.K.Huang, C.Kantner, R.C.O’Handley and V.V.Kokorin, Applied Physics Letter 69-13(1996)1966-1968 T.Kakeshita and K.Ullakko, MRS (Materials Research Society) Bulletin 27-2 (2002) 105 109 K.Inoue, Y. Yamaguchi, Y. Ishii, H. Yamauchi, and T. Shishido, Materials Science Forum 539- 543 (2007) pp. 3267-3272 A part of this work supported by “Nanotechnology Network Project of the Ministry of Education, Culture Sports, Science and Technology (MEXT), Japan” at the Research Center for Ultrahigh Voltage Electron Microscopy, Osaka University (Handai multi-functional Nano-Foundry).
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 463–464, DOI: 10.1007/978-3-540-85226-1_232, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1 a) SAED pattern taken from a direction (i.e., [010]T1 ) parallel to the interfacial plane. b) Schematic drawing of Bragg reflection spots in the reciprocal space.
Figure 2 a) DF image taken with spot 004T1 (T1) in Figure 1, and b) taken with spot 4 00T 2 (T2). c) Schematic drawing of the sample piece and the incident beam direction in real space.
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Nanocrystalline FeAl produced by high pressure torsion studied by TEM in 3D C. Mangler, C. Rentenberger and H.P. Karnthaler Physics of Nanostructured Materials, University of Vienna, Boltzmanngasse 5, 1090 Vienna, Austria [email protected] Keywords: severe plastic deformation, nanocrystalline, shear bands
High pressure torsion (HPT) is a very efficient way to carry out severe plastic deformation even on usually rather brittle intermetallics. Investigations on the structural evolution during HPT of different intermetallics revealed that in several cases the final homogeneous nanocrystalline structure is formed inhomogeneously via an intermediate structure consisting of nanocrystalline and coarse grained regions [1-3]. While investigations were carried out on top-view specimens, cross-sections were often neglected. In this work the structural evolution during the HPT deformation process was studied from top-view and cross-section using TEM and optical microscopy. FeAl single crystals were prepared from high purity metals by pre-alloying using the cold boat technique followed by Bridgman technique. All alloying procedures were performed under argon atmosphere. The single crystals were then deformed by HPT for up to 2 rotations using a pressure of 8GPa. Figure 1 shows TEM micrographs from a HPT deformed sample. To track the deformation process, the HPT disc was cut into two pieces along the rotation axis and the cut surface was then polished and etched. Investigations on the surface carried out by optical microscopy revealed a number of bands running from top to bottom – with increasing density towards higher radii. Cross section samples for TEM investigation were cut by FIB (cf. b and d in Figure 1a). Figure 1b shows a TEM bright field image of a cross-section prepared at a radius of about 1.5mm (corresponding to a shear deformation of ∼600%). A number of shear bands inclined with respect to the shear direction SD can be seen. With increasing deformation these shear bands broaden; they are nanocrystalline and disordered (cf. the top view in Fig.1c). From the comparison of the top-view and cross-section images it is concluded that the crystal is accommodating the shearing process by glide bands. The high activity in the glide bands acts as a trigger for the disordering. At larger radii (corresponding to higher grades of deformation) the bands are bent and stacked resulting in a rather layered nanocrystalline structure (cf. Fig.1d). With increasing deformation the grains become more equiaxed (cf. Fig.2). 1. 2. 3. 4.
C. Mangler, C. Rentenberger, H.P. Karnthaler, Proc. IMC16 (2006), p.1656. Rentenberger C., Mangler C., Karnthaler H.P. Mat. Sci. Eng. A 387- 389 (2004) 795-798. C. Rentenberger, H.P. Karnthaler Acta. Mater. 53 (2005) p. 3031 The authors thank Prof. Sassik, TU Wien for help during the initial alloying process, the group of Prof. Pippan for help with the HPT-deformation, Martina Dienstleder, TU Graz for the FIB preparation and the research project “Bulk Nanostructured Materials” within the research focus “Materials Science” of the University of Vienna.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 465–466, DOI: 10.1007/978-3-540-85226-1_233, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. Investigation of a HPT deformed sample. (a) Sketch of a HPT disk cut in half showing the locations where TEM samples (corresponding to (b) – (d) ) were taken; the shear direction is marked by SD. (b) TEM bright-field image of a cross section prepared by FIB from an area at a radius of about 1.5mm (deformation: 600%). Glidebands are visible. (c) Top-view image taken from another region showing disordered shear bands. The inserted SAD pattern was taken from the band. (d) TEM dark-field image taken at a radius of about 3mm (deformation: 1200%) showing a disordered nanocrystalline structure. Figure 2. TEM dark-field image, top view, of a sample deformed 9000%. The inserted SAD image shows a ring-pattern indicating a nanocrystalline structure. The absence of rings corresponding to super lattice shows the loss of long-range order.
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HRTEM study of precipitates in Al-Mg-Si-(Ag, Cu) alloys K. Matsuda1, J. Nakamura1, T. Kawabata1, T. Sato2, and S. Ikeno1 1. Graduate School of Science and Engineering for Research, University of Toyama, 3190, Gofuku, Toyama, 930-8555, Japan. 2. Graduate School of Science and Engineering, Tokyo Institute of Technology, 2-12-1, Ookayama, Keguro-ku, Tokyo, 152-8552, Japan. [email protected] Keywords: HRTEM, precipitate, crystal structure, chemical composition, aluminum alloys
It has been well known that excess Mg type Al-Mg-Si alloy causes better hardness, tensile strength and elongation than those of quasi-binary Al-Mg2Si alloy. Ag or Cu addition to Al-Mg-Si alloy also causes the modification of mechanical property, particularly its elongation [1,2]. We also had reported increase of tensile strength and uniform elongation in Ag or Cu addition into balanced Al-Mg2Si alloy. The precursor of Q’-phase, which was needle shape and its cross-section was elongated, existed in the sample aged at 523 K [3]. In the precipitation sequence of this alloy including Ag, we have found out the quaternary AlMgSiAg phase with hexagonal crystal system and a lattice parameter larger than that for the Q’-phase in Al-Mg-Si-Cu alloy [4]. Purpose of this study is to figure out an effect of the Ag or Cu addition on precipitates under the aging in excess Mg-type Al-1.0 mass% Mg2Si alloy using HRTEM. Al -1.0 mass% Mg2Si -0.4 mass% Mg (ex.Mg) alloys including 0.5 mass% Cu (ex.Mg-Cu) or 0.5 mass% Ag (ex. Mg-Ag) were used for the present work. The samples are hot- and cold-rolled till 1.0mm thickness. These samples were solution heat treated at 848K by atmospheric furnace for 3.6ks, quenched in ice water. Aging treatment was performed in the salt bath at 523K. AKASHI MVK-EII was used of hardness measurement by load of 100gf for 15s. Samples for TEM were polished by using two type of electrolyte, perchloric acid: ethanol=1:9 and nitric acid: methanol=1:3. HRTEM (Topcon EM-002B) with EDS was operated at 120kV and 200kV. 400kV HRTEM (JEOL4010T) was also used. Figure 1 shows age-hardening curves of three alloys aged at 523 K. The peak hardenss (HVmax)of ex.Mg-Cu alloy was the highest, the ex.Mg-Ag alloy is the second, and the ex.Mg alloy was the lowest among them because of the difference of amount of solute atoms. The time to HVmax was almost the same. Figure 2(a) and (b) are TEM pictures obtained for ex.Mg-Cu and ex.Mg-Ag alloys aged to HVmax(0.48ks). The fine needle-shaped precipitates were in there and no remarkable difference of microstructure between them. There were obvious difference from the alloy without Cu/Ag, however, has been found out in these precipitates. Precipitates of metastable phase in the ex.Mg alloy without Cu/Ag usually show a circle or an ellipsoidal shape of cross-section. Precipitates in alloys with Cu/Ag marked by arrows in Fig. 2(c) and (d) showed an elongated cross-section which was parallel to [010]Al or [100]Al. These are similar to the precursor of the Q’\phase in the balanced Al-Mg2Si alloy with Cu[3]. Details about
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HRTEM observation and analysis of crystal structure for these precipitates having elongated cross-section will be reported in the presentation. 1. 2. 3. 4.
K. Matsuda, T. Yoshida, T. Wada, Y. Uetani, T. Sato, A. Kamio, A. and S. Ikeno, J. Japan Inst. Metals, 62, (1998), p. 718. K. Matsuda, K. Kido, T. Kawabata, Y. Uetani and S. Ikeno, Journal of Japan Institute of Light Metals, 53 (2003), p. 528. K. Matsuda, D.Teguri, T. Sato, Y. Uetani, S. Ikeno, Materials Transactions, 48, (2007), p.967 K. Matsuda, S. Ikeno, T. Sato, Y. Uetani, Scripta Mater, 55, (2006), p.127.
Figure 1. Age-hardening curves obtained for alloys aged at 523 K.
Figure 2. TEM images obtained for alloys. (a) and (c): ex. Mg-Cu alloy, (b) and (d): ex.Mg-Ag alloy. Aging times for these pictures were 0.48 ks (HVmax) for (a) and (b), and 3.84 ks for (c) and (d).
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Martensite structure of non-stoichiometric Co2NiGa ferromagnetic shape memory alloy K. Prusik, M. Morawiec1 1. University of Silesia, Institute of Material Science, ul. Bankowa 12, 40-007 Katowice, Poland [email protected] Keywords: FSMA, martensite, γ-phase
A great deal of interest is attracted by alloys based on A2BC Heusler compounds with L21 superstructure. Until now the largest magnetic shape memory is exhibited by the alloys based on Ni2MnGa compound [1]. However, the practical application of these alloys in the polycrystalline state is limited because of their brittleness. In order to overcome the brittleness a new group of two-phase B2-γ Co-Ni-Ga alloys was developed as a system characterized by good ductility [2]. The B2 parent phase undergoes a martensitic transformation during cooling. Basing on the crystallographic model the structure of the martensite can be derived from the parent B2 lattice and described as the L10 (FCT) with the lattice parameters ratio c/a < 1 or BCT with c/a > 1. The studies were carried out on the two-phase martensite with the primary γ-phase particles alloys which chemical composition are given in Table I. Table I. Lattice parameters of the studied Co2NiGa alloys Alloy Co Ni Ga Martensite BCT at. % at. % at. % a [nm] c [nm] Alloy 1 48.6 23.3 28.1. 0.2703 0.3231 Alloy 2 49.0 21.2 29.8. 0.2711 03201 Alloy 3 49.3 22.1 28.6 0.2701 0.3221
γ phase a [nm] 0.3590 0.3592 0.3587
The structure of these alloys was studied both in the quenched at T = 1100oC and aged at 350oC state using HREM and X-ray diffraction methods. Experimental HREM images were obtained using JEOL JEM 3010 transmission electron microscope operating at 300 kV and attached with a CCD slowscan camera. The results show modulation of the martensitic structure. The extra spots in the diffraction patterns show inclination of the modulation vector of 13o from the [110] direction and its projection on the (001) plane is equal 0.68 nm (Figure 1). The HREM images, the corresponding FFT images and the extra spots in the diffraction pattern obtained for the martensitic structure of the alloy 2 with the excess of gallium atoms elucidate the role of these atoms in the lattice distortion. The lattice constants were determined by the X-ray method. The alloy ageing causes precipitation of fine dispersed coherent γ’ particles with L12 superstructure (Figure 2).
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1.
M. Richard, J. Feuchtwanger, D. Schlangel, T. Lograsso, S.M. Allen, R.C. O’Handley, Scripta Mater. 54 (2006) p. 1796 R.F. Hamilton, H. Sehitoglu, C. Efstathiou, H.J. Maier, Y. Chumlakov, X.Y. Zhang, Acta Mater. 54 (2006) 587
2.
(c)
(b)
(a)
(d)
[001]
Figure 1. Experimental HREM image for [001] orientation of the martensite with visible modulation (a) and enlarged area (d) with the corresponding FFT (b) and experimental diffraction pattern (c).
(a)
(b)
200
211
[011]
(c)
(d)
[001]
Figure 2. HREM images of the γ ’ precipitates for [011] and [001] orientation (a,c) and corresponding electron diffraction patterns (b,d)
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Electron microscopy of Fe and FeB atomic clusters in the Fe-based amorphous alloys structure E.V. Pustovalov, V.S. Plotnikov, B.N. Grudin, S.V. Dolzhikov, E.B. Modin, O.V. Voitenko and E.S. Slabzhennikov Electron microscopy Lab, Far Eastern National University, 8 Shukhanova str., Vladivostok, 690950 Russia [email protected] Keywords: electron microscopy, amorphous alloys, electron tomography
Amorphous Fe-base metallic alloys structure hierarchy were investigated by means of electron microscopy, structure modelling and tomography. Fe77Ni1Si9B13 and Fe67Ni6Si11B16 amorphous alloys structure were investigated. Atomic structures of amorphous FeB20 and FeB25 consists of 36000 atoms was modelled. For modelling and calculate electron tomography we use IMOD software [1]. Ordered structure regions, as a result of nonlinear image processing of HRTEM image, have various texture and represent “gratings”, “chains”, “clusters”. The mean periodicity (0.3 nm) in ordered regions and the effective size of a separate cluster (1-2 nm) were found. The symmetry of the ordered areas essentially differs from crystallographic because of their lattice sizes and, as a consequence, the arising considerable distortion in atomic structure [2, 3]. To clarify structure of the ordered regions we carry out series of calculating experiments of amorphous structure with embedded ordered atomic clusters. We modelled amorphous matrix of FeB20 and FeB25 using Ishikava approach with following structure relaxation [4]. Amorphous matrix models were consisting of 2500 – 36000 atoms. Ordered atomic clusters were prepared from crystals structure models α-Fe, FeB, Fe3B, Fe2B and Fe23B6. Unit cell was translated from 3 to 5 times and then some region was cut and transform to a new unit cell. Then prepared cluster was embedded into amorphous matrix (Fig.1). Such super cells have up to 30*30*10 nm dimensions and have from 2500 to 36000 atoms. Using multislice method we simulate HRTEM image series of tilted structure. Prepared images were used by IMOD for tomography reconstruction. There were various series of images were examined in the experiment. By results of tomogram reconstruction it is possible to make following conclusions: 1. For visualization of nanoclaster in an amorphous matrix its thickness should be proportional to thickness of a matrix. 2. It is necessary to use the series of images received under corners not less than-30 up to +30 degrees to an electronic beam. 3. As a result of work the tomogram of the α-Fe nonocluster in the amorphous matrix FeB20, shown on figure 3 has been reconstructed.
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4. For tomogram reconstruction we use single axis tomography method. As a result generated tomogram can’t give enough information for nanocluster three-dimensional modeling. In future experiments we plan to apply dual axis tomography method. 1. 2. 3. 4.
J.R. Kremer, D.N. Mastronarde and J.R. McIntosh, J. Struct. Biol. 116 (1996), p. 71. B. Grudin, S. Dolzhikov, V. Plotnokov, E. Pustovalov, Abstr. of Electron microscopy and multiscale modeling conf., Moscow, 3-7 Sept., p. 38. B. Grudin, S. Dolzhikov, V. Plotnokov, E. Pustovalov, Proc.of Asian Symp. Adv.Mater., Vladivostok, 1-4 Oct. 2007, p.O-14. T. Ichikawa, Phys.Stat.Sol. 29a (1975), p. 293. a
x y z
b c
Figure 1. Amorphous matrix structure FeB20 (black and gray) 3*3*1 nm with embedded α-Fe (lager violet) nanocluster. Size of Fe atoms in the nanocluster enlarged with better visibility.
Figure 2. Frame of the reconstructed tomogram α-Fe nanocluster in the amorphous matrix 9*9*3 nm.
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Core/Shell Precipitates in Al-Li-Sc-Zr Alloys V. Radmilovic1, M.D. Rossell1, A. Tolley1,2, E.A. Marquis3, R. Erni1 and U. Dahmen1 1. National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, University of California, One Cyclotron Rd., Berkeley, CA, 94720, USA 2. Centro Atomico Bariloche, CNEA, CONICET, 8400 San Carlos de Bariloche, Argentina 3. Department of Materials, Oxford University, Parks Road Oxford OX1 3PH UK [email protected] Keywords: Atom-probe-tomography (APT), electron microscopy, core/shell precipitates
Recent work on AlScZr alloys [1-3] has shown that a precipitate core-shell structure may be responsible for the modification of properties in a precipitate-hardened alloy system. It was demonstrated that the precipitates contained Sc and Zr, but that the distribution of Zr was not homogeneous. Instead, a Zr rich shell surrounded an Al3Sc core, acting as a diffusion barrier that reduced the growth rate of the Sc rich core. In this work we investigate the effect of Li addition to AlScZr alloys on core shell precipitate formation. The role of Li as a transient nucleating agent for Sc and Zr during heating and the formation of an Al3Li shell around (Sc,Zr)-rich particles during cooling is documented using atom probe tomography (APT) and energy-filtered electron microscopy (EFTEM), high angle annular dark field imaging (HAADF) as well as high resolution transmission electron microscopy (HREM). For the quaternary AlLiScZr alloys we show a way of producing a uniform distribution of monodisperse Al3(LiScZr) core/ shell inclusions in an Al matrix. Our approach uses differential diffusivities and solubilities of Li, Zr and Sc in an Al matrix. The dark field image in Figure 1a taken using a 110 superlattice reflection in an orientation near the [001] zone axis shows a typical distribution of monodisperse core/shell precipitates after heating for 18h at 450°C followed by 4h at 190°C. The precipitates exhibit a dark core surrounded by a brighter shell. The high angle annular dark field image in Figure 1b shows the contrast of the core region of the precipitates closely resembles that of the Al matrix. We show that for a core composition of Al3(Li.40Sc.48Al.12), measured by atom probe tomography (APT), the scattering power for the superlattice columns can nearly vanish even for a fully ordered L12 structure. As a result, the mean scattering intensity from the core is very similar to that of surrounding Al matrix.Figure 2 shows particles with a Li- and Sc-rich core surrounded by a Li-rich shell a few nanometers thick, with Zr segregated to core/shell interface. We propose the following precipitation mechanism: At low temperature, spinodal decomposition serves as a barrier-free process to grow evenly-spaced Li-rich clusters by congruent ordering [4], which act as heterogeneous nucleation sites for the formation of Sc-rich precipitates at high temperature. Once Li- and Sc-containing L12 precipitates are formed, Zr segregates at their surfaces during annealing at 450°C, while Li dissolves into Al forming Al rich solid solution. During the second stage annealing at 190ºC, the Li atoms form the shell around the Al3(LiSc) particles. This approach to generating
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precipitate distributions can be applied to a range of alloys and could lead to new types of dispersion-strengthened materials. 1. 2. 3. 4. 5.
A. Tolley, V. Radmilovic and U. Dahmen, Scripta Materialia 52 (2005) 621. C.B. Fuller and D.N. Seidman, Acta Materialia, 53 (2005) 5415. E. Clouet, L. Lae, T. Epicier, W. Lefebvre, M. Nastar, A. Deschamps, Nature Materials, 5 (2006)482. V. Radmilovic, A. Fox and G. Thomas, Acta Metallurgica, 37 (1989) 2385. This work was supported by the Director, Office of Energy Research, Office of Basic Energy Sciences, Materials Sciences Division, U. S. Department of Energy under Contract # DEAC02-05CH11231. HAADF
100 nm
1 nm
Figure 1. a) Dark field image of an Al-Li-Sc-Zr alloy aged 18h at 450°C and 4h at 190°C, recorded using 110 superlattice reflection; (b) HAADF image Al3(ScLiZr) core/shell precipitate.
Figure 2. 3D reconstruction (92nmx92nmx130nm) containing 3 precipitates and showing the Zr rich and Li rich regions around the Sc+Li rich cores.
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Analysis of basic mechanisms of hardening in ODS EUROFER97 steel using in-situ TEM A. Ramar1 and R. Schäublin1 1. Ecole Polytechnique Fédérale de Lausanne (EPFL), Center for Research in Plasma Physics, Association Euratom-Confédération Suisse, CH 5232 Villigen PSI, Switzerland [email protected] Keywords: Oxide, dislocation, yttria
Oxide dispersion strengthened (ODS) ferritic/martensitic EUROFER97 steel appears to be one of the promising candidates for high temperatures and nuclear applications such as structural material for the future fusion reactor. The presence of a fine dispersion of oxide particles enhances the inherent properties like swelling resistance and low radiation damage accumulation, of the base material EUROFER97 [1,2,3]. Recent studies have shown that the presence of nanometric Y2O3 particles in the matrix of EUROFER 97 improves its strength up to 600°C but there is a strong degradation in its mechanical properties at higher temperature [3,5]. The average size distribution of the oxide particles in the matrix is around 25 nm. The dispersed oxide, having higher strength than the EUROFER97 matrix, acts as a strong obstacle to the moving dislocations, which results in the alloy hardening. It was observed by the high-resolution TEM that the yttria particles are semi-coherent with the BCC EUROFER 97 matrix as shown in "Figure 1." This semi-coherency indicates that the shearing of yttria particles by dislocations would be difficult, leading to additional hardening. TEM observations were done on a JEOL 2010 microscope equipped with a LaB6 gun and a high tilt lens operated at 200 kV. Gatan TEM double tilt sample holder equipped with a molybdenum furnace operating up to 1000°C was used for in-situ heating TEM observations. In-situ TEM heating studies have shown that, upon heating, neither microstructure changes nor dislocation movement is observed up to 600°C. Movement of dislocations are observed above 680°C. Phase transformation to austenite starts at 840°C. Dislocation density decreases above 900°C. Yttria particles remains unchanged up to 1000°C as shown in “Figure 2.” Degradation in mechanical properties thus do not relate to changes in the yttria dispersion. To understand the type of interaction through which the dislocations overcomes the particle leading to the degradation in the mechanical properties above 500°C, in-situ TEM straining experiment were performed at room temperature and at 600°C. "Figure 3" shows, the dislocation pinning, bowing out and released from the particles during high temperature deformation. The interaction type could be an athermal or thermally activated mechanism, depending on the temperature. At temperatures below 400°C, athermal Orowan mechanism yields the main contributions to the flow stress [4]. At high temperatures, the activation energy for deformation calculated through variable strain rate experiments indicates that the dislocations overcome the particles through the
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thermally activated process of climb, however it is still remains to be seen [5]. Results obtained are presented here. 1. 2. 3. 4. 5.
D.K. Mukhopadhyay, F.H. Froes, D.S. Gelles, Journal of Nuclear Materials, 258-263 (1998), p. 1209. G.R. Romanoski, L.L. Snead, R.L. Klueh, D.T. Hoelzer, Journal of Nuclear Materials, 283287 (2000), p. 642. R. Schäublin, T. Leguey, P. Spätig, N. Baluc and M. Victoria, J. of Nucl. Mater. 307-311 (2002), p. 778. M. Bartscha, A. Wasilkowska, A. Czyrska-Filemonowicz and U. Messerschmidt, Mat. Sci and Engg. A, 272 (1999) p. 152 A.Ramar, P.Spätig and R. Schäublin, accepted for publication in J. of Nucl. Mater (2008).
Figure 1. (a) HR-TEM image of an yttria particle (b) on the matrix EUROFER 97, (c) FFT of the particles and (d) diffraction pattern taken in the matrix.
Figure 2. In-situ heating TEM bright field image of ODS EUROFER97 evolution with increasing temperature (arrows marks stable yttria particles)
Figure 3. TEM bright field image sequence of the dislocation interaction with an yttria particle (marked with a arrow) at 700°C
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TEM investigation on the acicular ferrite precipitation in γ’-Fe4N nitride X.-C. Xiong1, A. Redjaïmia1, M. Gouné1,2 1. Laboratoire de Science et Génie des Surfaces, URA-CNRS 1402, Ecole des Mines, Parc de Saurupt, 54042 Nancy Cedex, France 2. Automotive Products Research Center, ARCELOR RESEARCH, 57283 Maizièreslès-Metz Cedex, France [email protected] Keywords: Fe-N, pearlite, acicular ferrite
The Iron-Nitrogen system, similar to the Iron-Carbon, undergoes a eutectoid reaction which leads to the decomposition of austenite to ferrite and Fe4N nitride at the eutectoid point (592 °C and 2.4 wt.%). This eutectoid product, labelled nitrogen pearlite[1], is similar to the carbon pearlite. According to the Iron-Nitrogen phase diagram, the Fe4N nitride is nonstoichiometric and the nitrogen composition increases in Fe4N as the temperature decreases. This fact predicts the formation of ferrite in Fe4N phase when it is slowly cooled from the eutectoid temperature (592°C), i.e., γ ' → γ '+α . In this study, pure iron specimens were nitrided in austenite range and cooled slowly in the furnace in order to form nitrogen pearlite. The acicular ferrite has been found in the Fe4N nitride. Electron diffraction was used to determine the orientation relationship between the ferrite and the Fe4N nitride. The morphology of this acicular structure is clearly revealed by TEM (figure 1). Unlike the pearlitic structure in which all the lamellas (ferrite and Fe4N) grow in the same direction and are all parallel, the acicular crystals grow in Fe4N matrix in a way similar to Widmanstätten ferrite growing in an austenite grain. In order to identify the crystal structure of the acicular crystals, a series of TEM investigations has been conducted by applying a systematic microdiffraction method [1, 2]. From the I m3m space group (figure 2) and the lattice parameter a=0.287 nm measured by X-ray diffraction, this acicular crystal is concluded to be the α-ferrite. Both The Nishiyama-Wassermann (figure 3) and the Kurdyumov-Sachs (figure 4) orientation relationships [3,4] have been found between the acicular ferrite and the Fe4N matrix, i.e.:
1. 2. 3. 4.
( 011 ) α // (111 ) γ ' ; [100 ]α // [11 0 ]γ ' ; [ 011 ]α // [11 2 ]γ '
N-W OR
( 011 )α // (111 ) γ ' ; [11 1 ]α // [110 ]γ ' ; [ 211 ]α // [112 ]γ '
K-S OR
J. P. MORNIROLI and J. W. STEEDS, Ultramicroscopy 45 (1992) 219. A. REDJAIMIA and J. P. MORNIROLI, ibid. 53 (1994) 305. Z. NISHIYAMA, Sci. Rep. Tohoku Univ. 23 (1934) 638. G. WASSERMAN, Arch. Eisenhuttenwesen 16 (1933) 647.
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Fe4N matrix
Acicular ferrite
Proeutectoid ferrite
Figure 1. TEM image showing the acicular ferrite grown in the Fe4N matrix. γ’-Fe4N α-Ferrite
Zone axis : [100]α // [1-10]γ’
γ’-Fe4N α-Ferrite 1 α-Ferrite 2
Zone axis : [11-1]α//[1-10]γ’
7°
1-10α
220γ’
(011)α // (111)γ’
101α
011α 112γ’
222γ’ (-21-1)α // (112)γ'
002γ’ 002α
0-11α -1-12γ’
002γ’
-2-20γ’ 011α -2-22γ’ (011)α // (-1-11)γ’
(0-11)α // (-1-12)γ’
a
b
Figure 2. Electron diffraction pattern showing a) the Nishiyama-Wassermann and b) the Kurdyumov-Sachs orientation relationship between the acicular ferrite and the γ’-Fe4N matrix
479
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 A. Redjaïmia1, T. Otarola2 and A. Mateo2 1. Laboratoire de Science et Génie des Surfaces, Nancy-Université, CNRS. Parc de Saurupt CS 14234 F-54042 Nancy, France 2. ETSEIB – Universitat Politecnica de Catalunya Diagonal 647 08028 Barcelona [email protected] Keywords: Duplex (δ+γ) stainless steel, phase transformation, Orientation relationships, FrankKasper phases, Twinning, CSL and DSC lattices, Group theory.
Duplex stainless steels are being increasingly used as structural material in oil, chemical and power industries [1]. This is related to the fact that their duplex microstructure (δ+γ) allows a beneficial mixture of austenitic (γ) and ferritic (δ) properties: on the one hand high strength with a desirable toughness [2] and, on the other, good corrosion resistance. The ferritic phase (Im 3m, aδ = 0.2876 nm) , in the Fe-22Cr-5Ni-3Mo-0.03C ferriticaustenitic duplex stainless steel, can undergo a variety of decomposition processes when aged in the temperature range 650–550 °C. These processes are the precipitation of the twin related, γT, austenite particles, the Frank-Kasper χ and R phases. Both the crystal structure and the chemical composition of these phases have been studied in detail by electron microdiffraction and energy dispersive X-ray spectroscopy [3, 4]. These three phases γT, χ and R belong to (Fm3m, aγ = 0.359 nm) , (I 43m, a χ = 0.894 nm) and (R3, a R = bR = 1.07 nm; c R = 1.987 nm) space groups, respectively. The twin related, γT, austenite particles develop the Kurdjumov-Sachs (KS) orientation relationship (OR) with the ferritic matrix (Fig. 1a). Figure 2 points out 6 variants of austenite, which are the result of three twin related particles, leading to 24 variants in one ferritic grain. The χ Frank-Kasper phase develops a cube-cube OR with the ferritic matrix and a KS-OR with the twin related austenite variants (Fig. 1b). The Frank-Kasper R-phase develop the following OR (Fig. 1c): (0001)R // (111)δ ; [ 2110] R // [123] δ ; [1210] R //[ 312] δ with the ferritic matrix. The latter exhibits a KS OR with the twin related variants of the austenite (Fig. 1c): The derived triangular orientation relationships between the δ, γT and χ phases on one hand and between the δ, γT and R phases on the other hand, are analysed in terms of the symmetry theory [5] and of the coincidence site lattices (CSL and DSC) [6].
1. 2.
R. M. Davison and J . M. Redmond, Mater. Performance 29 (1990), p. 57. J. Foct, N. Akdut and G. Gottstein, Scripta Metall. Mater. 27 (1992), p. 1033.
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3. 4. 5. 6.
A. Redjaïmia, J.P. Morniroli, P. Donnadieu, G. Metauer, J. Mat. Sci. 37 (2002), p. 4079. A. Redjaïmia, A. Proult, P. Donnadieu, J.P. Morniroli, J. Mat. Sci. 39 (2004), p. 2371. J. W. Cahn and Kalonji, in Proceedings of the Conference on Solid-Solid Phase Transformations, ed. H. I. Aaronson, R. F. Seekereka, D. E. Laighlin and C. M. Waymann (Metal, Soc. AIME, Warrendale, (1982), p.3. W. Bollmann, Crystal Defects and Crystalline Interfaces, (1974) Springer, Berlin.
a) b) c) Figure 1. A series of electron diffraction patterns showing the orientation relationships between a) the δ-ferritic matrix and the twin related austenite recorded along δ// <110>γ; b) the δ-ferritic matrix, the twin related austenite and the χ-phase recorded along <111>δ//<110>γ//<111>χ; c) the δ-ferritic matrix, the twin related austenite and the R-phase recorded along <111>δ//<110>γ//[0001]R.
Figure 2. Composite electron diffraction pattern pointing out six variants (i.e., three twin related variants) of austenite exhibiting the Kurdjumov-Sachs orientation relationships with the δ-ferritic matrix. This diffraction pattern is recorded along δ // <110>γ.
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TEM study of localized deformation-induced disorder in intermetallic alloys of L12 structure C. Rentenberger, C. Mangler and H.P. Karnthaler Physics of Nanostructured Materials, Faculty of Physics, University of Vienna Boltzmanngasse 5, A-1090 Wien, Austria. [email protected] Keywords: severe plastic deformation, disordering of ordered structures, antiphase boundaries
Chemical atomic disorder of intermetallic alloys can be induced by severe plastic deformation (SPD). The process of disordering is usually monitored by the intensity of superlattice reflections using X-ray diffraction. In the present study transmission electron microscopy (TEM) methods are applied to study locally deformation-induced disordering in L12 ordered intermetallic alloys. Figure 1 shows a TEM bright-field image of a Ni3Al sample deformed by a shear strain of γ~2400%. As revealed from the selected area electron diffraction (SAD) patterns (cf. insets), a disordered nanocrystalline structure (NC) and an ordered coarsegrained structure (CG) are coexisting. The weak superlattice reflections (marked by arrows) in the SAD pattern indicate that in CG the order is reduced but still present. In the lattice plane image of CG, regions with different periodicity of the lattice fringes were encountered. In Figure 2 every other fringe shows a higher intensity as shown in the intensity profile. This agrees with image calculations based on L12 long-range order (cf. Figure 3). The observation of the superstructure of {100} lattice planes depends mainly on defocus and order. Slight variations of the orientation of the lattice planes with respect to the incoming beam do not change the intensity distribution considerably, whereas disordering leads to vanishing of the long-range periodicity (LRP) (cf. Figure 3b). Therefore, any indication of LRP is an evidence of residual chemical order and a reduction of LRP in an image can indicate defects reducing the order locally. Figure 4 shows a lattice plane image of a nanocrystal existing in region NC. Both LRP in the image and superlattice reflections in the corresponding power spectrum are missing. Figure 5 shows a lattice plane image of a nanocrystal taken under the same imaging condition as in Figure 4. Contrary, a residual superlattice structure can be detected locally in the corresponding power spectrum. The occurrence of some residual chemical order after SPD is consistent with a model of deformation-induced disorder based on the formation of antiphase boundaries (APB) [2]. Since at APBs many next neighbours are chemically disordered they are lattice defects reducing the order locally. In a nutshell, TEM methods are very useful for studying deformation-induced disorder. The TEM results support a process of disorder by the formation of APB faults. 1. 2. 3.
C. Rentenberger and H. P. Karnthaler, Acta mater 53 (2005), p. 3031. C. Rentenberger and H.P. Karnthaler, Acta mater (2008) in press. We kindly acknowledge the support by the “Bulk Nanostructured Materials” research project within the “Materials Science” research focus of the University of Vienna.
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Figure 1. Ni3Al. TEM bright-field image and selected area diffraction patterns of the duplex structure consisting of disordered nanocrystalline (NC) and of ordered coarse grained (CG) structure.
Figure 2. Ni3Al. TEM lattice plane image of the coarse grained area showing the long-range periodic structure. Simulated image and intensity profile perpendicular to the fringes are shown as insets. Figure 3. Ni3Al. Simulated (a) contrast of lattice fringes of a 5nm thick ordered (a) or disordered (b) sample under different defocus Δf (b) (-60/-80/-100 nm) conditions (beam direction ~[016]).
Figure 4. Ni3Al. TEM lattice fringe image of a nanocrystal showing {200} lattice planes; in the corresponding power spectrum fundamental {200} reflections are present only.
Figure 5. Ni3Al. TEM lattice fringe image of a nanocrystal and the corresponding power spectrum (superlattice SL and fundamental F reflection) indicate some local residual order.
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SEM and TEM study of dynamic recrystallisation of zirconium alloy L. Saintoyant1,2, L. Legras1 and Y. Brechet2 1. EDF R&D, Les Renardières, Département MMC, 77818 Moret sur Loing, France 2. SIMAP, INPGrenoble, CNRS UJF, Domaine Universitaire, BP 75, 38402 Saint Martin d’Hères cedex, France [email protected] Keywords: recrystallization, creep, zirconium
Fuel claddings of pressurised water reactor are made of zirconium alloys. They are the primary barriers between fuel pellets and the external environment. During service, claddings are irradiated and oxidised as well as hydrided because of the release of hydrogen during corrosion reaction. During transportation of spent fuel, temperature and stresses transients lead to the evolution of the microstructure of the fuel cladding. Recovery of the pre-existing dislocation introduced during cold working and even sometimes recrystallisation may occur. This work intends to study the microstructural evolution under such conditions through orientation mapping using Transmission Electron Microscope (TEM) FEI TECNAI F20 G2 equipped with ACT device, a Scanning Electron Microscope (SEM) FEG LEO Supra 55 and a dual beam FEI Helios Nanolab both equipped with Electron Backscattered Diffraction (EBSD) and also 3D EBSD system Autoreveal in the case of the dualbeam. An experimental study on the effect of creep on recrystallisation is carried out. Interrupted creep test of zircalloy-4 sheet are performed at 470°C with stresses of 80MPa, 100MPa and 120MPa. The microstructural characterisation is first carried out by TEM. Figure 1 presents the evolution of the dislocation structure during creep. Figures 1a and 1b show that after a 15h creep test, recrystallisation takes place. It is underline by the formation of nucleus in highly deformed area and at grains boundaries. Nevertheless, the initial grain structure can still be seen. If the duration of creep tests increases, formation of cells (1c) and, at the end, a partially recrystallised microstructure is observed (1d) erasing the initial cold work elongated grain. At the same time, hardness tests show that this recrystallisation process is accelerated by the applied stresses. To identify the proper mechanisms undergoing during the dynamic recrystallisation further TEM characterisations are carried out. The effect of disorientation between grains or cells is investigated. Indeed, Zircalloy-4 has a strong deformation and recrystallisation texture [1-2]. The microtexture is therefore explored thanks to EBSD. Figure 2 presents the EBSD map obtained after a creep test at 470°C during 110h. Similarly to figure 1d, recrystallised and deformed areas can to be seen. But, due to the limits of the EBSD technique, deformed areas are poorly indexed. Consequently, microtexture in deformed area is also measured in TEM, thanks to the ACT software developed by Dingley and Wright [3], which uses reconstructed diffraction patterns from conical dark fields.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 483–484, DOI: 10.1007/978-3-540-85226-1_242, © Springer-Verlag Berlin Heidelberg 2008
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Finally 3B EBSD is performed to determine the localisation and distribution of nuclei. Grains boundaries are analysed to understand the effect of Coincidence site lattice orientation [1] on recristallisation. 1. 2. 3.
K. Zhu in “Etude des mécanismes de deformation et de recristallisation dans un alliage de zirconium”, PhD of the University Paris XIII, France (2006). F. Gerspach in « Mécanismes d’évolution de texture au cours du recuit d’alliages de zirconium et de titane », PhD of the University Paul Verlaine, France (2007). S.I. Wright and D.J. Dingley, Materials Science Forum 273-275 (1998), p. 209.
Figure 1. TEM micrograph of the microstructure of crept zircalloy-4 at 470°C after different solicitation time; (a) and (b) 15h, εt=1,56%, (c) 24h, εt=2,8%, (d) 110h,εt=9,2%. Scale bar is 500nm for (a) to (c) and 5μm for (d).
Figure 2. EBSD orientation map of zircalloy-4 deformed to 9,2% at 470°C. Scale bar is 5μm.
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Effects of solution treatment and test temperature on tensile properties of high strength high Mn austenitic steels K. Phiu-on1, W. Bleck1, A. Schwedt2 and J. Mayer2 1. Institut für Eisenhüttenkunde, RWTH Aachen, Germany 2. Gemeinschaftslabor für Elektronenmikroskopie, RWTH Aachen, Germany [email protected] Keywords: high manganese austenitic steel, solution treatment, deformation twinning, TWIP, TRIP, EBSD
Tensile properties of high strength high Mn austenitic Fe-26.5%Mn-3.6%Al2.2%Si-0.38%C-0.005%B and Fe-18.9%Mn-0.62%C-0.02%Ti-0.005%B-0.11%N steels were investigated after different solution treatments. The results show that the solution treatment has a significant influence on microstructure and mechanical properties of the investigated steels. By appropriate solution treatment the product of tensile strength (Rm) and total elongation (A50) of the hot rolled steel can be improved from 40,00050,000 MPa.% to 55,000-65,000 MPa.% depending on the steel chemical composition. In addition, tensile tests were carried out at different temperatures in order to analyse the correlation between deformation mechanisms and corresponding mechanical properties. The microstructure of the steels without deformation and after deformation by tensile test was investigated by EBSD. For the range of investigation, reducing the initial grain size before deformation does not change the main deformation mechanism, but it has a significant influence on mechanical properties of the investigated steels. A solution treatment with a rather high temperature, e.g. 1100 °C for the Fe-18.9%Mn-0.62%C-0.02%Ti-0.005%B-0.11%N steel, results in a significant increase in the ε-martensite fraction during quenching. This deteriorates the ductility of the steel. The EBSD measurements reveal the mechanisms contributing to the overall plasticity of the investigated steels on the microscale. The plasticity of the 26.5%Mn3.6%Al-2.2%Si-0.38%C-0.005%B steel is produced mainly by TWIP (twinning induced plasticity) mechanisms under the examined experimental conditions, whereas for the Fe-18.9%Mn-0.62%C-0.02%Ti-0.005%B-0.11%N steel TWIP and TRIP (transformation induced plasticity) mechanisms occur with different degrees depending on the test temperature of the tensile test. For the contributing phase transformations the EBSD measurements point to the transformation sequence austenite → ε-martensite → α’-martensite. 1.
We gratefully acknowledge the financial support of the Deutsche Forschungsgemeinschaft (DFG) within the Collaborative Research Centre (SFB) 761 “Stahl – ab initio. Quantenmechanisch geführtes Design neuer Eisenbasis-Werkstoffe”.
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Figure 1. EBSD results of the Fe-18.9%Mn-0.62%C-0.02%Ti-0.005%B-0.11%N steel after hot rolling, 1000°C/30min solution treatment and tensile test at room temperature. From left to right: Image quality map including Σ3 twin boundaries in the austenite partition marked in yellow, phase map, IPF maps of the sample normal direction for the single phases (Confidence index > 0.1) .
Figure 2. EBSD results of the Fe-18.9%Mn-0.62%C-0.02%Ti-0.005%B-0.11%N steel after hot rolling, 1000°C/30min solution treatment and tensile test at -40°C. Image sequence and legend as in Fig. 1.
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Microstructure evolution during Ni/Al multilayer reactions S. Simões1, F. Viana1, A.S. Ramos2, M.T. Vieira2 and M.F. Vieira1 1. Dep. de Engenharia Metalúrgica e Materiais, GMM/IMAT, Faculdade de Engenharia, Universidade do Porto, Rua Dr. Roberto Frias, 4200–465 Porto, Portugal 2. ICEMS, Dep. de Engenharia Mecânica, Faculdade de Ciências e Tecnologia, Universidade de Coimbra, R. Luís Reis Santos, 3030-788 Coimbra, Portugal [email protected] Keywords: nickel aluminides, multilayer, exothermic reaction
Heating of thin multilayer systems can lead to stable or metastable phase formation. The structural evolution is known to depend on processing pathways and/or nominal composition; interdiffusion or intermixing prior to phase transformation, play a key role in determining the phase formation kinetics. The reaction of Ni and Al multilayer to form NiAl has been the subject of some investigations over the past two decades. The first phase to form depends on sample processing, bilayer thickness and annealing conditions. In the literature the first phase to form has been reported to be Al3Ni [1], AlNi [2], or Al9Ni2 [3]. Ni/Al multilayer is an interesting system since it is known to transform to NiAl in a rapid, exothermic, self sustained reaction. The heat is released locally and progress through the multilayer until all the Al and Ni has reacted. This system has potential for use as heat source for joining temperature-sensitive materials, like microelectronic components. Thin films with nanometric Ni and Al alternated layers were deposited by d.c. magnetron sputtering using nickel and aluminium pure targets. In our experiments, we use periods, or bilayer thicknesses, of 5, 14 and 30 nm and a total film thickness ranging from 2 to 5 μm. The structural evolution of the Ni/Al multilayer with temperature was studied by differential scanning calorimetry (DSC), scanning electron microscope (SEM), transmission electron microscopy (TEM) and X-ray diffraction (XRD). DSC experiments were performed on freestanding films, from room temperature to 700 °C at 10 and 40 °C/min. Two exothermic reactions were detected in DSC curves of the film with 30 nm period, with peak temperatures at 230 and 330 ºC while for the 5 and 14 nm period film we only observed one exothermic peak at 190 and 250ºC, respectively. To identify the reaction products, DSC samples were examined by XRD. For the asdeposited condition only Al and Ni were detected by XRD. The films with 30 nm period were heated at 300ºC (temperature between the two peaks), 450ºC (temperature after the second peak) and 700ºC. For the 300ºC sample, the XRD patterns identified Al3Ni and Ni phases. For the 450ºC sample, Al3Ni and NiAl were the detected phases. Only after heating up to 700ºC, Ni and Al react completely to form NiAl. For the 5 and 14 nm periods films, the formation of NiAl occurs in one single step. The structural evolution was followed by SEM and TEM to observe the morphological changes occurring during reaction. Figure 1 shows SEM images for as-deposited and DSC samples treated at 300ºC and 700ºC. With the formation of NiAl the multilayer morphology is no longer present. S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 487–488, DOI: 10.1007/978-3-540-85226-1_244, © Springer-Verlag Berlin Heidelberg 2008
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1. 2. 3.
E. Ma, C.V. Thompson, L.A. Clevenger, J Appl Phys 69 (1991) p. 2211 C. Michaelsen, G. Lucadamo, K. Barmak, J Appl Phys 80 (1996) p. 6689. M.H. Silva Bassani, J.H. Perepezko, A.S. Edelstein, R.K. Everett, Scripta Mat 37 (1997) p. 227.
This work was supported by “Fundação para a Ciência e a Tecnologia” through the project PTDC/CTM/69645/2006 and the Grant SFRH/BD/30371/2006 financed by POS_C. a
b
c
NiAl
Figure 1. Schematic illustration and SEM images of the structural evolution of Ni/Al multilayer samples: a) as-deposited, b) annealed at 300ºC and c) annealed at 700ºC.
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TEM investigation of severely deformed NiTi and NiTiHf shape memory alloys G. Steiner, M. Peterlechner, T. Waitz and H.P. Karnthaler Physics of Nanostructured Materials, University of Vienna, Boltzmanngasse 5, 1090 Vienna, Austria [email protected] Keywords: TEM, amorphization, nanocrystalline, NiTi, NiTiHf
NiTi shape memory alloys show a martensitic phase transformation from a cubic high temperature phase (B2 austenite) to a monoclinic structure (B19´ martensite). A crystalline to amorphous phase transformation can be obtained by methods of severe plastic deformation such as cold rolling and high pressure torsion (HPT) [1]. Deformation at a temperature below the martensitic finish temperature (Mf) promotes the amorphization [2]. A NiTi alloy and a NiTiHf alloy that are martensitic at room temperature (RT) but have different values of Mf (~30 and 110°C, resp.) were subjected to HPT at RT and analysed by transmission electron microscopy (TEM). Discs (8mm ø, 0.8 mm thick) were deformed by HPT (4 GPa, 12 turns). TEM specimens were punched out at a distance of 2.7 mm from the centre of the HPT discs corresponding to a deformation of 250. TEM was carried out at 200 kV. The phase structure was analysed by selected area diffraction pattern (SADP) methods. The TEM bright field image of Fig. 1a shows the heterogeneous microstructure of NiTi obtained by HPT. Band shaped areas of highly strained and fragmented martensitic grains (cf. A in Fig. 1a and the corresponding SADP in Fig. 1b). Martensite was observed containing nanoscale (001) compound twins; these deformation twins could facilitate the amorphization acting as obstacles and causing dislocation accumulation [3]. Amorphous phase occurs near A since in Fig. 1b diffuse rings superimpose the diffraction spots of the crystalline lattice. Areas that contain a mixture of a nanocrystalline and an amorphous phase are observed near B in Fig. 1a (cf. the corresponding SADP of Fig. 1c showing a ring pattern containing diffraction spots of the nanocrystals and broad diffuse rings of the amorphous phase). Diffraction rings were observed that correspond to B2 and B19´. Therefore, in some of the nanograins a reverse transformation from the martensite to the austenite was induced during the HPT. Compared to NiTi, considerably less nanocrystalline and amorphous phase is observed in NiTiHf after HPT (cf. Fig. 2). B2 austenite is not observed in SADP of NiTiHf. NiTiHf contains rather large grains (> 300 nm) with relatively few dislocations leading to moderate lattice strains only. A stripe like contrast is frequently observed that might arise by a twinned martensitic lattice. The present results indicate that dynamical recovery hinders dislocation accumulation and therefore the formation of a nanocrystalline structure and the transition to an amorphous phase. 1. 2.
T. Waitz, V. Kazykhanov, H.P. Karnthaler, Acta Mater. 52 (2005) p. 137. S.D. Prokoshkin et al., Acta Mater. 53 (2005) p. 2703.
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3. 4.
H. Nakayama et al., Mater. Trans. 42 (2001) p. 1987. The authors thank the group of Prof. R. Pippan for the kind help with the HPT deformation. Support by the research project "Bulk Nanostructured Materials" within the research focus "Materials Science" of the University of Vienna and the I.K. "Experimental Materials Science – Nanostructured Materials" a college for Ph.D. students is acknowledged.
Figure 1. HPT NiTi (a) TEM bright field image showing highly deformed and fragmented grains (near A) and a mixture of nanograins and amorphous phase (near B). (b) and (c) SADP of the areas marked by A and B, resp., in (a).
Figure 2. HPT NiTiHf. (a) TEM bright field image of twinned grains containing relatively few dislocations. (b) SADP of the area encircled in (a) showing diffraction spots and weak diffuse rings caused by a small volume fraction of amorphous phase.
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TEM studies of nanostructured NiTiCo shape memory alloy for medical applications D. Stróż and Z. Lekston Institute of Material Science, University of Silesia, Bankowa 12, 40-007 Katowice Poland [email protected] Keywords: shape memory effects, nanocrystalline microstructure, interface structure
The NiTi shape memory alloys have been widely studied over the last 50 years as they show the best functional properties and find applications in many different fields of industry and medicine. Recently, seeking for still improved properties of these alloys, the attentions was focussed on the opportunity of controlling the alloy structure by adjusting their crystallisation conditions. The alloys are subjected to severe plastic deformation - leading to their amorphization - and then annealed at relatively low temperatures that produces nanocrystalline structure of the required grain size [1-4]. It was found that if the grains size is less than 60 nm only the R phase transformation can occur in the alloy, while the grains larger then 150 nm contained the B19’ martensite only that showed a unique “herring-bone” structure [3]. In the recent work the binary NiTi alloy was modified by addition of 1.3 at.% of Co substituting nickel that improved the alloy workability [5]. It was designed for producing implants in form of clamps used in surgical treatment of mandibular fractures. In order to ensure the recovery temperatures close to the human body temperature and optimise its functional properties, the alloy was subjected to different thermo-mechanical treatment, one of which was cold-rolling and annealing. Samples in form of 2 mm wire were cold rolled by 30% and then annealed at the temperature range 350oC – 600oC. It was found that annealing at 350oC and 400oC produced nanocrystalline structure where the grain sizes varied from about 50 – 200 nm (Figure 1). The specimens showed a very good superelasticity effect caused by the R-phase transformation. The B19’ martensite was not formed in these samples – as could be seen at the DSC curves. The TEM observation carried out at room temperature showed mainly the B2 phase in the sample, occasionally in same grains the R-phase was found. The HREM studies of the grain boundaries structure revealed that many of them were coherent twin boundaries (Figure 2). Often small angle boundaries were also observed. This could be the reason of the good functional properties of these samples. 1. 2. 3. 4. 5.
C. Rentenberger, T. Waitz, H.P. Karnthaler, Scripta Materialia 51 (2004), p.789 T.Waitz, H.P.Karnthaler, Acta Materialia 52 (2004), p. 5461 T. Waitz, Acta Materialia 53 (2005), p. 2273 X.Wang, J.J. Vlassak, Scripta Materialia 54 (2006), p. 925 J. Drugacz, Z. Lekston, H. Morawiec, K. Januszewski, J. Oral Maxillofac.Sur. 53 (1995), p. 665
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a
b
Figure 1. Nanocrystalline structure of the cold-rolled and then annealed at 350oC/1h NiTiCo alloy – bright field (a) and dark field (b) images
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b
c
Figure 2. High resolution image of a grain boundary between the nanograins (a) and the processed image (b) showing the twin relationship between both grains obtained by filtering the FFT (c)
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TEM investigations of microalloyed steels with Nb, V and Ti after different treatments G. Szalay1, R. Grill2, K. Spiradek-Hahn1, and M. Brabetz1 1. Austrian Research Centers GmbH- ARC, Advanced Materials and Aerospace Technology, Alloy Development Group, 2444 Seibersdorf, Austria 2. voestalpine Grobblech GmbH, Voestalpine-Str. 3, 4020 Linz, Austria [email protected] Keywords: low carbon steel, heat treatment, carbides
The low carbon steel microalloyed with Nb, V, and Ti has been thermomechanically rolled with subsequent accelerated cooling (TM+DIC). This results in a homogeneous fine grain microstructure which leads to high strength and ductility [1]. Additionally a part of this material has been quenched and tempered (TM+DIC+QT). For the analysis of the microstructure a FEI CM-20 STEM Transmissionelectronmicroscope at an acceleration voltage of 200kV was used. The TEM is equipped with a secondary electron detector (SE) for scanning the foil surface and a energy dispersive X-ray spectrometer (EDX) for chemical analysis. For the analysis of the crystal structure of the precipitates electron diffraction was used. The chemical composition of the particles was analysed by EDX. HRTEM investigations on ultra-fine precipitates were carried out with a FEI Tecnai F20 at 200kV. The initial state (TM+DIC) exhibits fine long and narrow grains as a result of thermomechanical rolling with a lathlike substructure inside the grains. At the grain boundaries and lath interfaces a very low density of cementite precipitates was observed (cf. Fig 1a). Layers of retained austenite were present at the lath interfaces (cf. Fig 1b). Inside the laths a low density of cementite (50-80nm in size) and near cubic TiC or (Ti,Nb)C precipitates (50-100nm in size) were inhomogeneous distributed (cf. Fig 1c). Contrary to this the additionally quenched and tempered steel (TM+DIC+QT) shows a high density of homogeneous distributed precipitates at laths interfaces and grain boundaries (cf. Fig 2a) identified by electron diffraction as cementite. No more retained austenite at lath interface has been observed. Inside the laths coarse cementite (50300nm) and (Ti,Nb) carbides (50-100nm), with near cubic morphology and fine (<25nm) round precipitates were homogeneous distributed (cf. Fig 2b). EDX analysis and HRTEM confirms the ultrafine precipitates (<15nm) as mainly Nb2C and (Nb,Mo) carbides, respectively (cf. Fig 2c) So additional quenching and tempering leads to a high density of homogeneous distributed cementite precipitates at grain boundaries and inside the laths. Also a high density of fine (7-25nm) (Nb,Mo,V) carbides can be observed inside the laths. Furthermore the retained austenite was no more observed and dislocation density was significantly reduced. 1 H.K.D.H. Bhadeshia, R.W.K. Honeycombe, “Steels-Microstructure and Properties”, (Butterworth-Heinemann) (2006), p. 214.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 493–494, DOI: 10.1007/978-3-540-85226-1_247, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1 a) SE image of TEM foil, microstructure overview b) Bright field, dark field images of retained austenite at the lath interfaces c) Bright field images of cementite precipitates (left) and a TiC particle (right) inside a lath
Figure 2 a) SE image of TEM foil, microstructure overview b) Bright field images of (Ti, Nb) carbides smaller than 25nm c) HRTEM image, EDX spectrum and FFT spectrum of a ultra-fine Nb2C particle
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Initial Stages of the ω Phase Transformation R. Tewari, K.V. Manikrishna, G.K. Dey, S. Banerjee Materials Science Division, Bhabha Atomic Research Centre, Mumbai, India 400 085 [email protected] Keywords: phase transformation, diffuse intensity, simulation
The ω phase was initially thought to be a phase encountered only in group IV elements and its based alloys. Recent results, however, have shown that the formation of the phase is not confined to group IV but it is an ‘inbuilt transformation’ into any metallic system containing body centre cubic crystal structure. One of the least understood part of the β→ω transformation is the pre-transition effects associated with the transformation. The β→ω transformation shows pronounced pre-transition effects in terms of the appearance of diffuse intensity in diffraction patterns, phonon softening along certain wave vectors of the bcc lattice, anomalous diffusion, etc [1,2]. The appearance of diffuse scattering prior to the β→ω transition has prompted intense theoretical and experimental activities [1, 2]. One of the most noticeable features in the distribution of the diffuse intensity is shift in the maxima of diffuse intensity to incommensurate positions from the ideal ω position in the reciprocal space (Fig.1). Several models, in literature, have been proposed to predict the extent of the shift in the incommensurate positions of the diffuse maxima [3-5]. Most of these models, however, met with partial success. In the present paper, the formation of ω phase under electron irradiation and by thermal treatment has been examined in a Zr-20 %Nb alloy. The parameters which have been varied during irradiation were the temperature and the time of irradiation. Thermal treatments as well as irradiation were found to lead to the formation of diffuse omega maxima. The origin of these maxima was probed in either case by carrying out high resolution electron microscopy (HREM). A comparison of omega formation by these treatments has shown subtle differences. Attempts were made to explain the differences by simulating the initial stages of the ω phase transformation using Monte Carlo methods. The model uses different structural configurations of the embryos of ω to generate the SAD patterns with desired shift in the diffuse intensity and correlated with the internal structure of the ω-embryo. The predicted dimensions of these ω-embryo were found in close agreement with those observed in the high resolution images 1. 2. 3. 4. 5.
S.K. Sikka, Y. K. Vohra and R. Chidambaram, Prog. Mater. Sci.,27, (1982) 245. C.Stassis, J. Zarestky, N.Wakabayashi; Phys. Rev. Lett., 41, (1978) 1726. D. deFontaine and R. Kikuchi, Acta Metall., 22, (1974) 1139. R. Strychor, J. C. Williams and W. A. Soffa, Metalls. Trans., 19 A (1988), 226. H. Ezaki, M. Morinaga, M. Kato and N. Yukawa, Acta Metall. Mater.,39, (1991), 1755-1761 W. Sinkler and D. E. Luzzi, Acta Metall. Mater., 42, (1994) 1249.
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Figure 1. Distribution of diffuse intensity in different selected area diffraction patterns of the bcc lattice.
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TEM study of the Ni-Ti shape memory micro-wire H. Tian1, D. Schryvers1, and J. Van Humbeeck2 1 EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium 2 MTM, K.U.Leuven, Kasteelpark Arenberg 44, B-3001 Leuven, Belgium [email protected] Keywords: Micro-wire, TEM, Surface layers, ELNES, Compositional gradient
Today, due to its unique mechanical properties, durability and biocompatibility, applications of Nitinol in a wide variety of medical implants are progressively increasing [1-3]. However, as Nitinol consists of about 50 at. % Ni, certain applications are still hindered by the concern of free Ni release in the surrounding tissue. It is known that the biocompatibility of the implants made from TiNi alloys depends on a corrosion resistant titanium oxide layer avoiding the severe toxicological and allergic responses of Ni. The latter is controlled by the structure of (near-) surface layers which can strongly be affected by various surface treatments [3,4]. In this present work, the oxide surface of an as-received Nitinol micro-wire is investigated by different TEM techniques. FIB was used to prepare near surface cross-section sample. Spectroscopical quantification results using EDX and EELS reveal that the main contents of the surface oxide layer are Ti and O. In order to improve the precision and to obtain better spatial localization of the EELS data, an analysis with nano-probe was performed. Nano-probe ELNES identified the native oxide layer at the surface as a combination of TiO and TiO2 (rutile), although some Ni containing particles are still observed (Figure 1). Rutile and anatase TiO2 can be differentiated by a small difference in the shape of the second peak of Ti L3 edge [5,6]. The interior of the wire is divided into a core-shell structure with primarily B2 grains in the shell, which is approximately 1 micron thick, and heavily twinned B19’ martensite in the central core, as show in Figure 2. In between the metallic wire and the oxide surface, an interfacial nanolayer containing Ni3Ti particles is observed. The B19’-B2 core-shell structure is the result of an exponentional variation in the chemical composition in the metal wire which is revealed by both EDX and EELS as well as EFTEM (Figure 3). The observed interface between austenite and martensite perfectly fits with the expected composition at room temperature and the compositional gradient is related to the formation of the Ni3Ti intermediate layer and the oxide surface. 1. 2. 3. 4. 5. 6.
S. Kujila, A. Pajala, M. Kallioinen, et al., 2004, Biomaterials 25, 353-358 J. Ryhänen, 1999, Biocompatibility Evaluation of Nickel-Titanium Shape Memory Metal Alloy, Ph.D. Thesis, Oulu University, Oulu, Finland S. A. Shabalovskaya, 2002, Bio-Medical Materials and Engineering 12, 69-109 S. A. Shabalovskaya, 1996, Bio-Medical Materials and Engineering 6, 267-289 R. Brydson, H. Sauer, W. Engel, 1991, TMS Annual Meeting, New Orleans, p. 131. P. Potapov, W. Tirry, D. Schryvers, V.G.M. Sivel, M.-Y. Wu, D. Aslanidis, H. Zandbergen, 2007, J. Mater. Sci.: Materials for Medicine 18, 483-492
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Figure 1. Nano-probe ELNES of the Ti L3,2 edges in different regions of oxide layer
Figure 2. (a)Bright field image overview of the near surface and bulk structure; (b) HRTEM of nanoscale twin structures of B19' martensite in the core.
Figure 3. Averaged Ni content in the Ni-Ti microwire close to the surface as measured from the profile of the Ni/Ti ratio map plus fitted exponential function (the Ni3Ti layer starts about 100nm further to the right of the last data point).
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Multi-scale observations of deformation twins in Ti6Al4V W. Tirry1, F. Coghe1, L. Rabet1 and D. Schryvers2 1. COBO, Royal Military Academy, Renaissancelaan 30, 1000 Brussels, Belgium 2. EMAT, Universiteit Antwerpen, Groenenborgerlaan 171, 2020 Antwerp, Belgium [email protected] Keywords: Ti6Al4V, deformation twin, TEM
Deformation twinning is an important and necessary mechanism to accommodate deformations in materials with a hexagonal close packed (hcp) crystal structure, i.e. commercial pure titanium. Ti6Al4V is a dual phase alloy of which the major phase (+90%) is the hcp alpha-phase. The occurrence of deformation twins in this type of material is not well documented yet, especially not in the case of large deformations like the shear deformations presented here. This work presents a multi-scale analysis of deformation twins present after a simple shear deformation applied on Ti6Al4V plate material. The applied shear strains are 10%, 30% and 50%. A combination of optical microscopy, electron backscatter diffraction (EBSD), conventional and high resolution transmission electron microscopy are used to study the twin type, density and morphology. Optical microscopy was used to have a first impression on whether twins are present or not, with the advantage of having a wide field of view. After optimising the metallographic preparation, a fraction of alpha grains in the 30% and 50% samples shows optical contrasts typical for twinning. A CM20 and FEG equipped CM30 transmission electron microscope were used to obtain conventional and high resolution observations of the deformation twins. Using conventional TEM all twins were identified as being of the {10-12} type. Figure 1.a shows a BF image of a grain in which two orientation variants of the {10-12} twin are present. The inclination of the twinning plane trace with respect to the macroscopic shear direction can be approximately determined and seems always to be almost parallel or almost perpendicular to each other. Diffraction patterns obtained from regions inside the deformation twin show streaking along the [0001] direction, as illustrated in Figure 1.b. This streaking is not present in the original and the residual parent grain. BF and DF images of the twins interior structure reveal a high density of fringes and streaks perpendicular to the [0001] direction, indicating the possible presence of multiple stacking faults on the basal plane. This phenomenon was already observed in deformation twins induced in Zr and commercially pure titanium [1]. Figure 2 shows a HRTEM observation of the interior structure of a {10-12} twin. Figure 2.b is an enlarged selection revealing an ABABACACA stacking sequence inside the twin. Furthermore optical and TEM observations already revealed twinning is not occurring in all grains. As such, EBSD measurements and orientation microscopy are used to obtain more statistical data on the occurrence of twins and this in relation to the crystallographic texture and applied shear direction.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 499–500, DOI: 10.1007/978-3-540-85226-1_250, © Springer-Verlag Berlin Heidelberg 2008
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S.G. Song and G.T. Gray III, Phil. Mag. A 71, no. 2 (1995), p. 263. Acknowledgement: This work was carried out in the frame of the "Inter University Attraction Poles program-Belgian Science Policy" under contract number IUAP P6-27.
Figure 1..a) Two variants of {10-12} twins, the white arrow approximately indicates the macroscopical shear direction b) DP of area 1 and 2 in a.
Figure 2. a) HRTEM image of a {10-12} twin showing internal stacking faults as indicated by the white square and magnified in b
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Nd:YAG laser joining between stainless steel and nickel-titanium shape memory alloys J. Vannod1,2, A. Hessler-Wyser1 and M. Rappaz2 1. CIME, Ecole Polytechnique Fédérale (EPFL), 1015 Lausanne, Switzerland 2. LSMX, Ecole Polytechnique Fédérale (EPFL), 1015 Lausanne, Switzerland [email protected] Keywords: laser welding, SMA, stainless steel, diffusion couple, SEM
Nickel-Titanium (NiTi) shape memory alloys are often used in medical component devices, for instance as guide wires for neurological surgery applications [1]. The manufacture of such devices becomes more and more challenging, especially considering the need to join them with other metals. Laser welding is a promising technique to realize and to guaranty the mechanical stability of dissimilar metal welds, although inherent differences in chemical compositions, absorption, physical and thermo-mechanical properties can lead to severe problems, in particular fracture of the weld due to the possible formation of brittle intermediate phases. Laser welds of NiTi - stainless steel (SS) pieces have been made with a Nd:YAG laser and the weld microstructure (see Figure 1) ha been studied by Scanning Electron Microscopy (SEM). The phases and defects in these welds have been compared with those observed in autogenous NiTi welds and SS welds (see Figure 2). They have been put into relation with the Fe-Ni-Ti ternary phase diagram [2]. In addition, Differential Thermal Analyses (DTA) of NiTi-SS alloys and NiTi-SS diffusion couple experiments have been performed in order to gain a better understanding of the phases and reactions occurring during laser welding. This diffusive couple experiments have been analyzed by Energy Dispersive X-ray Spectroscopy (EDX) It appeared that although experiments were performed under controlled atomosphere, oxide layers have restrained the chemical diffusion of concerned elements (see Figure 3). Further diffusion couple experiments will be realized with stainless steel welded caps to avoid oxygen contamination during heating. 1. 2.
T. Duerig, A. Pelton, and D. Stöckel, “An overview of Nitinol medical applications” Materials Science and Engineering A, 273-275, (1999), p. 149-160. G. Cacciamani, J. De Keyzer, R. Ferro, U. Klotz, J. Lacaze, and P. Wollants, “Critical evaluation of the Fe-Ni, Fe-Ti and Fe-Ni-Ti alloy systems”, Intermetallics, 14, (2006), p. 1312-1325.
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Figure 1. Scanning Electron Micrograph of bimetallic laser welds between NickelTitanium SMA (Top) and Stainless Steel 304 (Bottom).
Figure 2. Scanning Electron Micrograph of autogenous laser welds in Nickel-Titanium SMA (Left) and Stainless Steel 304 (Right).
Figure 3. Scanning Electron Micrograph of diffusion couple between NiTi SMA (Top) and Stainless Steel 304 (Bottom). The presence of Rutile (TiO2), in between, is due to oxygen contamination during the experiment and restrains other element’s diffusion.
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Focused Ion Beam application on the investigation of tungsten-based materials for fusion application L. Veleva1, R. Schäublin1, A. Ramar1, Z. Oksiuta1, N. Baluc1 1.Ecole Polytechnique Fédérale de Lausanne (EPFL), Centre de Recherches en Physique des Plasmas, 5232Villigen PSI, Switzerland [email protected] Keywords: FIB, SEM, tungsten, powder metallurgy
Tungsten-base materials are candidate materials for structural applications in the high temperature region of plasma facing components for future fusion power reactors (DEMO and beyond) [1, 2]. Of promising interest are oxide dispersion strengthened (ODS) tungsten-based materials, which allow improving the mechanical properties of the base material, in particular the creep strength at high temperatures. It is also expected that such materials exhibit improved resistance to radiation damage, as the numerous interfaces between the particles and the matrix may act as sinks for the irradiation-induced defects. In our study we have used two different ways to produce ODS W alloys, where the oxide consists in yttria (Y2O3), by additions to pure tungsten of either pure yttria or yttrium. Mechanical alloying (MA) is being used to produce intimate mixing of these materials [3]. This method allows obtaining a homogeneous distribution of oxides inside tungsten. Unfortunately, after compaction under 200 MPa and sintering at 2000°C our material is porous and brittle. While the porosity amounts to 10% up to 15% of the volume, which explains the brittleness, it appears that the dislocation density is relatively low, which would allow for fair ductility. In order to assess the microstructure (oxide dispersion, precipitation, dislocations), and ultimately to improve on the production of these ODS tungsten alloys, SEM and TEM observations are performed on the compacted material. However, the high porosity makes the preparation of transmission electron microscopy (TEM) specimens in a classical way unreliable and further microstructural investigation difficult if not impossible. Focused ion beam (FIB) appears then as a major tool in this study, as it allows milling and observing specimens in any region of the material. We have used for FIB/SEM operation a ZEISS NVision 40 located at PSI. With the help of the lift-out FIB technique, we were able to prepare TEM lamellae that are 10x5 μm2 and 150 nm thick from the fragile, porous and fine-grained microstructure of the tungsten-based materials. However, TEM sample preparation by FIB introduces radiation damage in the surface of the thin lamella that will interfere with the radiation damage we would like to study in the bulk of the material. In order to remove the radiation damage induced by the FIB in the surfaces, we have developed a flash electrochemical polishing technique that is applied after the FIB preparation. FIB was also used to study the porosity of the materials. Using the ‘slice and view’ technique, it was possible to create a 3D reconstruction of the bulk and study the type of porosity in the W-Y and W- Y2O3 materials (Figs 1a and 2a respectively). It appears that porosity is closed in the W-Y
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 503–504, DOI: 10.1007/978-3-540-85226-1_252, © Springer-Verlag Berlin Heidelberg 2008
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material (Fig. 1b), showing spherical pores, while it is percolating in W-Y2O3 material (Fig. 2b), showing connecting wide empty channels. Results will be presented here. 1. 2. 3.
R. Neu, V. Bobkov, R. Dux, A. Kallenbach, Journal of Nuclear Materials 363-365 (2007) 52. N. Baluc, Final Report on the EFDA Task TW1-TTMA-002_05. M.-N. Avettand-Fenoel, R. Taillard, J. Dhers, J. Foct, International Journal of Refractory Materials and Hard Materials 21 (2003) 205-213.
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Figure 1. FIB/SEM 3D reconstruction of W-Y material after sintering using the in-lens backscattered electron detector showing (a) the material and (b) the corresponding porosity consisting in isolated cavities (e.g. arrow). Acquisition time: 4 hours with 27 nA 30 keV FIB beam.
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Figure 2. FIB/SEM 3D reconstruction of W-Y2O3 material after sintering using the inlens backscattered electron detector showing (a) the material and (b) the corresponding percolating porosity consisting in empty and wide connected channels. Acquisition time: 4 hours with 27 nA 30 keV FIB beam. ☺
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HRTEM of NiTi shape memory alloys made amorphous by high pressure torsion T. Waitz1, K. Tsuchiya2, M. Peterlechner1 and H.P. Karnthaler1 1. Physics of Nanostructured Materials, University of Vienna, Boltzmanngasse 5, 1090 Vienna, Austria 2. Innovative Materials Research Laboratory, National Institute for Materials Science, Sengen1-2-1, Tsukuba, Ibaraki 305-0047, Japan [email protected] Keywords: HRTEM, amorphization, NiTi
NiTi shape memory alloys show a martensitic phase transformation from the ordered cubic B2 high temperature phase (austenite) to the monoclinic B19´ martensite. Amorphization of bulk NiTi can be obtained by severe plastic deformation using high pressure torsion (HPT) [1]. Crystalline debris surviving the HPT triggers nanocrystallization that occurs upon annealing of amorphous HPT NiTi [2]. Nanograins impact the martensitic phase transformation [3] and nanocrystalline alloys can be obtained that show enhanced shape memory and superelasticity. In the present paper, high resolution transmission electron microscopy (HRTEM) was applied to investigate the debris structure embedded in the amorphous matrix. NiTi shape memory alloys that are martensitic at room temperature were subjected to high pressure torsion to obtain a crystalline to amorphous phase transformation. Specimens were punched out from the HPT discs and used to prepare TEM foils by twin jet polishing. HRTEM was carried out using a CM 30ST operating at 300 kV. Fig. 1 shows a HRTEM micrograph of the amorphous matrix. After a strain of 1300, the crystalline structure is almost completely destroyed. However, in HRTEM images areas 1-3 nm in size are observed that that show atomic correlations similar to that of a crystalline lattice. This medium range ordered debris shows a rather homogeneous distribution. Embedded in the amorphous matrix also larger nanocrystals (< 20 nm) are left over after the HPT deformation (cf. Figs. 2 and 3). These nanocrystals are heterogeneously distributed since frequently clusters of them are observed. Fig. 2a shows a nanocrystal that is about 9 nm in size. The nanocrystal contains B2 austenite (cf. the power spectrum of Fig. 2b); spots are observed that correspond to {001} superlattice reflections. Therefore, chemical order is still encountered in the nanocrystals that have survived the severe plastic deformation. Fig. 3a shows a HRTEM micrograph of a nanocrystal of ~17 nm size. In the power spectrum, spots arise that correspond to a B19´ martensitic lattice structure. B19´ nanocrystals that have a size less than about 15 nm were not observed since smaller crystals are B2 only. This size dependent phase stability of nanocrystals surviving HPT is similar to that observed in the case of thermally induced martensitic phase transformations in NiTi nanograins [3]. 1.
T. Waitz, V. Kazykhanov, H.P. Karnthaler, Acta Mater. 52 (2005) p. 137.
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M. Peterlechner, T. Waitz and H.P. Karnthaler, Proc. Intern. Symp. Bulk Nanostr. Mater., Ufa, Russia (2007) p. 57. T. Waitz, T. Antretter, F.D. Fischer, N.K. Simha, H.P. Karnthaler, J. Mech. Phys. Sol. 55 (2007) p. 419. Support by the University of Vienna (research focus "Materials Science" and I.K. "Experimental Materials Science – Nanostructured Materials") is acknowledged.
Figure 1. HPT NiTi. HRTEM micrograph of the amorphous phase. Areas of medium range order 1-3 nm in size are indicated.
(b) (a) Figure 2. HPT NiTi. (a) HRTEM micrograph of a B2 nanocrystal embedded in the amorphous matrix. (b) Power spectrum. {001} superlattice reflections are indicated.
(b) (a) Figure 3. HPT NiTi. (a) HRTEM micrograph of a martensitic nanocrystal embedded in the amorphous matrix. (b) Power spectrum corresponding to a [10 1] zone axis of B19´.
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Is the lattice structure of the martensite in nanocrystalline NiTi base centred orthorhombic? T. Waitz Physics of Nanostructured Materials, University of Vienna, Boltzmanngasse 5, 1090 Vienna, Austria [email protected] Keywords: HRTEM, martensite, NiTi
NiTi shape memory alloys show a martensitic phase transformation from the ordered cubic B2 lattice (austenite) to a low temperature phase called martensite. Experimentally, the martensitic lattice structure of single crystalline and coarse grained NiTi was found to be monoclinic (B19´, space group P21/m) [1]. In striking contrast, using ab-initio total energy methods the ground state of NiTi was calculated to be base centred orthorhombic (BCO, space group Cmcm) [2]. Recently, it was argued that atomic scale (001) compound twinning observed in nanocrystalline NiTi is related to the BCO structure [3]. In the present paper, high resolution transmission electron microscopy (HRTEM) was applied to investigate the lattice structure of the atomic scale compound twins in nanocrystalline NiTi. Nanocrystalline NiTi was obtained via devitrification of an intermediate amorphous phase. Applying different beam directions (BD), HRTEM was carried out using a CM 30ST operating at 300 kV. Fig. 1a shows a sketch of the B19´ lattice (monoclinic angle β ≈ 98°). The sketch of Fig. 1b shows the unit cell of BCO. BCO can be regarded as a sequence of unit cells of B19´ (each containing 2 (002) planes) that are (001) twin related and have a specific value of β = 107.2°. Fig. 2 shows an HRTEM image of atomically twinned martensite observed in nanocrystalline NiTi. The smallest twins have a width of 4 (002) planes (thickness of 0.9 nm) i.e. they comprise 2 B19´ unit cells. This thickness is twice the value expected in the case of the BCO lattice (cf. Fig. 1b). Using BD = [010], the monoclinic angle can be measured directly in Fig. 2; the result of 6-8° agrees with the value expected in the case of B19´. Similar, in the case of BD = [110] the comparison of experimental HRTEM images (cf. Fig. 3a) with image simulations shows good agreement with B19´ but disagrees with BCO (cf. Figs. 3b and c, resp.). It is therefore concluded that in nanocrystalline NiTi the martensite is B19´ having a monoclinic angle similar to that observed in coarse martensitic grains of NiTi. Since the smallest twins contain 2 unit cells of B19´ instead of one unit cell only, experimental evidence for a relation of the atomic scale twinning to the BCO structure is not obtained. The occurrence of the atomic scale (001) compound twins in nanocrystalline B19´ NiTi is facilitated by their very low specific twin boundary energy and enables the accommodation of transformation strains in the nanograins [4].
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K. Otsuka, X. Ren, Progr. Mater. Sci. 50 (2005) p.211. X. Huang, G.J. Ackland and K.M. Rabe, Nature Mater. 2 (2003) p. 307. J. Zhao, F.L. Meng, W.T. Zheng, A. Li and Q. Jiang, Mater. Lett. 62 (2008) p. 964. T. Waitz, T. Antretter, F.D. Fischer, N.K. Simha, H.P. Karnthaler, J. Mech. Phys. Sol. 55 (2007) p. 419. Support by the University of Vienna (research focus "Materials Science") is acknowledged.
(a) (b) Figure 1. NiTi. Sketch of different atomic structures of the martensite. (Ni and Ti are indicated by squares and circles; full and open symbols indicate atoms of alternating (020) layers). (a) Unit cell of B19´ (β ≈ 98°). (b) BCO. The unit cell (marked by dashed lines) contains two twin related units of B19´ (indicated by full lines, β = 107.2°).
Figure 2. Nanocrystalline NiTi. HRTEM micrograph of (001) compound twins. Twin boundaries and (101) lattice planes are marked by dashed and full lines, resp. The numbers indicate the thickness of the twins in units of (002) lattice planes. (BD=[010]).
(a) (b) (c) Figure 3. NiTi. HRTEM images (BD= [110] ). (a) Experimental HRTEM image. (b) Simulated HRTEM image based on B19´. (c) Simulated HRTEM image based on BCO.
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TEM study of the NiTi shape memory thin film B. Wang1, A. Safi2, T. Pardoen2, A. Boe3, J.P. Raskin3, X. Wang4, J.J. Vlassak4 and D. Schryvers1 1. EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020, Antwerp, Belgium 2. Department of Materials Science and Processes, Université catholique de Louvain (UCL), IMAP, Place Sainte Barbe 2, B-1348 Louvain-la-Neuve, Belgium 3. Department of Electrical Engineering, ELEC, Université catholique de Louvain (UCL), B-1348 Louvain-la-Neuve, Belgium 4. Division of Engineering and Applied Sciences, Harvard University, Cambridge, MA 02138-2901 [email protected] Keywords: thin film, FIB, TEM, columnar grains
Recently, NiTi thin films have received a growing interest owing to the large stress induced transformation strain and the high work density involved in that transformation [1]. Crystalline NiTi thin films made by sputtering are known to exhibit superelastic behaviour and robust shape memory effect applicable when resilient mechanical response is essential, or when thermal actuation or sensing is desired [2]. In the present work, a TiNi thin film was deposited on the following stack: Si3N4 / Si / Si3N4 / TiNi (80nm/700μm/80nm/700nm). The presence of a thin Si3N4 sacrificial layer allows the release of freestanding films. The thin film is fully crystallized at 460 ℃ and the composition corresponds to 51at% Ti. An advanced cross-section sample preparation technique using focused ion beam (FIB) is employed for the TEM sample preparation. The sample is extracted from the NiTi surface perpendicular to the stack. Preliminary TEM observation shows that the film is inhomogeneous with a columnar grain morphology mainly showing the B2 structure (Figure. 1). Also, SAED patterns obtained from some grains reveal the presence of the R-phase. Pt layer and Au layer deposited prior to FIB milling aimed at protecting the NiTi thin film. The BF image shows that the columnar grains are inhomogeneous with a typical diameter equal to about 1 μm. Also some precipitates, which could be Ti2Ni, are observed in the film. The columnar grain orientation inside the thin film is not random, but involves several directions. In the near future, freestanding films undergoing different levels of mechanical loading will be prepared directly for TEM investigation, by removing the thin Si3N4 sacrificial layer. The actuation will be provided by an actuator beam or layer involving large internal stresses (Figure. 2) which, upon release, will pull on the NiTi sample. 1. 2. 3. 4.
K.S.S. Eswar Raju, S. Bysakh, M.A. Sumesh, et al., Mater. Sci. Eng. A 476 (2008), 267-273 Michale J. Vestel, David S. Grummon, Mater. Sci. Eng. A 378 (2004), 437-442 D. Fabrègue, N. Andre, M. Coulombier et al., Micro & Nano Letters 2 (2007), 13-16 We kindly acknowledge support of the projects FWO G.0465.05 and IAP P6/24
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Figure 1. BF image of the NiTi thin film with Si substrate
Figure 2. Schematic illustration of removing Si3N4 sacrificial layer [3]
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Sub-nano analysis of fine complex carbide in high strength steel with probe Cs (S)TEM K. Yamada1, E. Hamada1, K. Sato1 and K. Inoke2 1. Steel research Laboratory, JFE steel corp. 2. Japan FEI [email protected] Keywords: carbide, high tensile strength steel, sub-nano analysis
Distribution of fine carbides is quite effective to achieve high tensile strength and excellent formability simultaneously in hot rolled steels. Recently, high performance steels are strongly required as materials for automobile parts because using high strength steel make possible to reduce size and weight. Attaining lighter car bodies has a great advantage from the point of view of improving fuel efficiency and reduction of emission of carbon dioxides. In order to obtain high strength in such steel, the control of fine carbides of several nm in diameter is a key technology in industrial steel processing. [1] Although nano analysis using conventional FE-TEM has been widely used in the steel research field for better understanding precipitation in practical steel, the interests of researchers is now shifting to the sub-nano field such as the interface between individual carbides and matrices as well as the atomic structure of grain boundaries. Higher analytical potential with Cs corrected STEM is now thought to be powerful to clarify the sub-nano structure of various carbides. In this study, we focused in the precipitation of fine complex carbides in hot-rolled steel with the tensile strength 800MPa. Specimens were prepared with a conventional twin-jet type electro-polishing technique using an electrolytic solution (mixture of perchloric acid, methanol, and 2-n Butoxyethanol) at 253K. High-resolution TEM and STEM experiments were conducted using a TITAN (80-300) with a Cs probe corrector. Figure 1 shows a typical TEM micrograph of (Ti,Mo)-C carbides that has a specific distribution feature called inter-phase precipitation. As shown in this Fig.1, individual carbides show a platelet or needle-like shape. This is categorized as an MC type with a NaCl structure. The crystallographic relationship between the carbide and the bcc matrix is like the Baker-Nutting's one without two of the three equivalent variants as reported by A.T. Davenport in V added steels. [2] Figure 2 shows an HAADF STEM image of a similar area of Fig.1 and an example of a STEM/EDX line analysis across an approximately 5nm carbide using an approximately 0.3nm electron probe. In this figure, the contrast of carbide is higher than that of the ferrite matrix but this is mainly due to a local change in thickness. The atomic ratio of Ti and Mo in this carbide seems to be constant and this agrees with the former result using conventional FE-TEM. [3] However, a more accurate analytical condition should be selected because this carbide has a specific thin shape. So, we are planning to do the same analysis under accurate edge-on conditions of the target carbide.
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Y.Funakawa, T.Shiozaki, K.Tomita, T.Yamamoto and E. Maeda: ISIJ International, 44(2004), 1945-1951 2. K.Sato, H.Nakamichi and K. Yamada: Kenbikyou, 40, No.3 (2005), 183-187. 3. A.T.Davenport, F.G. Berry and R.W.K. Honeycombe: Met. Sci. J. 2 (1968), 104.
2nm Figure 1. High resolution TEM micrograph showing inter-phase precipitation area.
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Figure 2. HAADF STEM image of carbides and their typical STEM/EDX line analysis across a single carbide indicated by the arrow.
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Characterization of morphology and microstructure of different kinds of materials at NTNU Mater Sci EM Lab Y.D. Yu, T. Nilsen, M.P. Raanes, J. Hjelen and J.K. Solberg Department of Materials Science and Engineering, Norwegian University of Science and Technology (NTNU), N-7491 Trondheim Norway [email protected] Keywords: EPMA, SEM, TEM
The NTNU Materials Science and Engineering EM Lab is equipped with the electron microscopy facilities assisting NTNU multi-disciplinary materials science research. The laboratory services research activities from various material departments at the university. Here we present several recent characterization examples from different kinds of materials. Figure 1(a) shows a digital transmission electron microscopy (TEM) micrograph of an Al-Mg-Si alloy from our JEM-2010 TEM, which is equipped with a sensitive CCD camera. Further semi-automatic particle analysis by using imaging software could give the detailed information of the hardening phase distribution in this commercial Al alloy, this distribution being as an important parameter for understanding the microstructure– mechanical property relationship. The development of bulk nanostructured material (BNM) from this commercial alloy produced by severe plastic deformation (SPD) [1] has been revealed at atomic level by high resolution TEM as shown in Fig. 1(b). Catalyst particle characterization is the basis for catalyst production and the performance of such particles. Figure 2(a) shows a secondary electron equivalent image from our Zeiss scanning electron microscope (SEM), equipped with an in-lens detector (Zeiss Supra 55VP). In this case a careful sample preparation was employed to disperse particles exactly at the same imaging plane for archiving nanosacle resolution at a short working distance (WD). The high-resolution TEM image of Fig. 2(b) confirmed the full crystallization of the individual particles, completing the full range of microstructure characterization of this catalysis system [2]. Figure 3 shows a variable pressure (VP) SEM micrograph of a nano-composite polymer, in which the SEM gas pressure was carefully controlled to form the thinnest conducting ionizing layer for getting the best VP SEM resolution. Without using conducting coating, the nano-enhanced particles in this polymer system have definitely been identified in the VP SEM mode [3]. Furthermore, precise quantitative element measurements and X-ray mapping can be performed by our electron probe micro analyser (EPMA) JXA-8500F, and Figure 4 shows a submicron resolution X-ray mapping from a pig iron sample where small oxide inclusions could be identified inside large titanium-nitride lamellar particles. The EM Lab is also involved in intensive research programmes for electron back scattering diffraction (EBSD), and some results from this work are also presented at the conference. 1. 2.
M. Liu, H.J. Roven and Y.D. Yu, Zeitschrift für Metallkunde 3 (2007), p. 184. T.J. Zhao, D. Chen et al., Norwegian Patent Application (2008).
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L. Shao, PhD Thesis (in press), Norwegian University of Science and Technology (2008).
Figure 1a. High-resolution TEM image of the hardening phase distribution in a commercial Al-Mg-Si alloy.
Figure 1b. A screw dislocation dissociated into two Shockley partials connected by an intrinsic stacking fault in SPD aluminium.
Figure 2a. SEM micrograph of catalyst particle distribution at WD of 4 mm.
Figure 2b. High-resolution TEM image of one of the catalyst particles in Fig. 2(a).
Figure 3. VP SEM micrograph of a nano composite polymer at the 8 Pa gas pressure.
Figure 4. X-ray mapping from a pig iron, showing the inclusions of oxides within lamellar nitrides.
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Characterization of a Ti64Ni20Pd16 thin film by transmission electron microscopy R. Zarnetta1,2, E. Zelaya2,3, G. Eggeler2 and A. Ludwig1,2 1. Combinatorial Material Science group, Caesar, Ludwig-Erhard-Allee 2, 53175 Bonn, Germany. 2. Ruhr-University Bochum, Institute of Materials, 44780 Bochum, Germany. 3. Centro Atómico Bariloche, Av. Bustillo 9500, 8400 Bariloche-RN, Argentina. [email protected] Keywords: thin film characterization, Ti-Ni-Pd, shape memory alloys, TEM, XEDS.
Ti–Ni-thin films fabricated by sputter-deposition are applicable in microdevices like: micropumps, microwrappers, and microgrippers [1–5], since they exhibit excellent shape memory effects and good mechanical properties [6]. The addition of a third element like Pd, Pt or Hf increases the transformation temperature. It is already reported that the substitution of Pd for Ni can increase the transformation temperatures of Ti-Ni from approximately room temperature to 800 K [7]. The phase transformation characteristics within the Ti-Ni-Pd system were recently investigated by combinatorial materials science methods [8]. For the first time the complete composition region showing a reversible phase transformation was determined. For non-transforming compositions, e.g. Ti64Ni20Pd16, X-ray diffraction results showed pronounced amounts of Ti2Pd and Ti2Ni precipitates. However the mechanism of how the precipitates might hinder the phase transformation could not be deduced. Therefore a TEM sample of Ti64Ni20Pd16 was cut form the Ti-Ni-Pd composition spread by target preparation methods using a focused ion beam (FIB) system with attached energy dispersive X-ray (EDX) analysis. Using TEM, two different types of phases could be distinguished within the sample using composition determination by XEDS and microdiffraction pattern (MDP). Fig. 1a shows a bright field (BF) image of the thin film and phases with distinctively different microstructures. Some grains exhibit directional striation and others a flat contrast. Also some rounded-like precipitates are pointed with arrows. The composition of the different microstructures was determined by XEDS technique. In figure 1b a R-X line spectrum through the grains with different microstructure and a precipitate is shown (dash white line in figure 1b). Depletion in Pd content is observed on the grains with a striated microstructure and in the precipitate region. Fig. 2a shows a MDP of the phase exhibiting a striated contrast within the BF image. This pattern seems to correspond to a compact phase structure. The determination of the structure is still under investigation. Fig. 2b shows a characteristic MDP of the phase with flat contrast. The spots of this pattern could be indexed corresponding to a [111] zone axis of the B2 structure. From obtained results it can be concluded that two different types of precipitates are present. Additionally a new compact phase was found that may additionally inhibit the martensitic transformation of the present B2 phase within the Ti64Ni20Pd16 sample. S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 515–516, DOI: 10.1007/978-3-540-85226-1_258, © Springer-Verlag Berlin Heidelberg 2008
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1. 2. 3. 4. 5. 6. 7. 8. 9.
D.D. Shin, K.P. Mohanchandra, G.P. Carman, Sensor Actuator A 111 (2004) 166. E. Makino, T. Mitsuya, T. Shibata, Sensor Actuator A 88 (2001) 256. J. Gill, D. Chang, L. Momoda, G. Carman, Sensor Actuator A 93 (2001) 148. M. Tabib-Azar, B. Sutapun, M. Huff, Sensor Actuator 77 (1999) 34. M. Kohl, D. Dittmann, E. Quandt, B. Winzek, S. Miyazaki, D.M. Allen, Mater Sci Eng A 273-275 (1999) 784. S. Miyazaki, A. Ishida, Mater Sci Eng A 273–275 (1999) 106. J. Wu, Q. Tian, Intermetallics 11 (2003) 773. R. Zarnetta, A. Savan, S. Thienhaus, A. Ludwig, App Surf Sci 254 (3) (2007) 743. Y.C. Lo, S. K. Wu, Scripta Metall Materialia 27 (1992) 1097.
Figure 1. a) Bright field image of a Ti64Ni20Pd16 thin film, b) EDX line scan along the dashed line indicated in the left micrograph.
Figure 2. Microdiffraction patterns of the grains a) with a striated microstructure, b) with a flat contrast. The left pattern was taken with a nominal spot size value of 40 nm and the right one with a nominal spot size value of 13 nm.
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Analytical electron microscopy investigations of a microstructure of single and polycrystalline β-Mg2Al3 Samson phase A. Zielińska-Lipiec, B. Dubiel and A. Czyrska-Filemonowicz Faculty of Metals Engineering and Industrial Computer Science, AGH University of Science and Technology, Al. Mickiewicza 30, 30-059 Kraków, [email protected] Keywords: Complex Metallic Alloys, β-Al3Mg2, Samson phase
The β-phase (Samson phase) of the Al-Mg system, with large unit cells containing 1168 atoms per unit cell, belongs to the class of the Complex Metallic Alloys (CMAs). The β-Mg2Al3 has a very low density and unique combinations of properties which are mutually excluded in conventional materials [1,2]. Two single crystals (as-grown and annealed at 180˚C/53h) and one as-grown polycrystalline specimens of the β-Mg2Al3 were investigated by XRD, LM as well as analytical SEM and TEM. TEM investigations were performed by JEM-2010ARP microscope using thin foils prepared by ion beam milling. The electron diffraction paterns (SAED) were interpreted by the JEMS software. The microstructure of as-grown single crystal specimen of composition Mg38.5Al61.5 was characterized by an inhomogeneous (striated) contrast of a matrix due to strain field (Fig.1). After annealing at 180ºC for 53 h, an increased number of planar defects (domain boundaries) in two perpendicular <220> directions was observed (Fig.2). In a matrix two phases were coexisting: the cubic β-phase (Fd-3mS: ICSD 57964, Samson phase) and rhombohedral β’-phase (R3m). The β-Mg2Al3 unit cell contents 1168 atoms while the β’-Mg2Al3 293 atoms only (model II, hR293) [1,3,4]. SEM investigations of a polycrystalline specimen revealed large grains of Samson phase with small precipitates enriched with Mg (probably γ-phase). TEM analysis showed the presence of lamellar grains with a small misalignment to each other (Fig.3a). Inside the grains, sparsely distributed planar defects (domain boundaries) and dissociated dislocations were observed (Figs 3b-d). Electron diffraction and XRD analyses confirmed that the microstructure of the investigated specimen was consisting of the β-Mg2Al3 Samson phase. The β’-phase was not detected. 1. 2. 3. 4. 5.
S. Samson, Acta Crystallogr., 19 (1965), p. 401 M. Feuerbacher et al., Zeitschrift für Kristallographie 222 (2007), p. 259 Q.B.Yang, Andersson, L. Stenberg, Acta Crystallogr. B43 (1987), p. 14 J.Timm, H. A. Warlimont, Zeitschrift für Metallkunde 71 (1980), p. 434 The authors acknowledge a financial support from the European Network of Excellence (NoE) Complex Metallic Alloys, contract no NHP-CT-2005-500140. Valuable contribution of Prof. Philippe A. Buffat (EPFL) to phase identification is kindly acknowledged.
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Figure 1. Matrix striated contrast in as- Figure 2. Planar defects in the single grown single crystal sample. crystal sample annealed at 180˚C for 53h. a)
b)
c)
d)
Figure 3. Microstructure of a polycrystalline β-Mg2Al3 specimen: a) lamellar subgrains, b) planar defects and dislocations, c) plate-like grain and striated contrast in the adjacent grains, d) dissociated dislocations.
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Grain boundary interfaces in ceramics D.J.H. Cockayne, S.-J. Shih, K. Dudeck and N. Young Department of Materials, University of Oxford, Parks Road, Oxford OX1 3PH, UK [email protected] Keywords: interfaces, high resolution
Grain boundary interfaces play an important role in the microstructural properties and sintering behaviour of ceramics. Of particular interest is the presence and nature of intergranular films (IGF), the planarity, termination and roughness of the interfaces, and segregation. The resolution of the electron microscope is such that these questions are all amenable to study, since, in thin crystals, the positions of individual atomic columns or even individual atoms can be determined. The goal is to compare experimentally determined structures with those obtained from atomistic models, with a view to giving confidence to predictive modelling. However obtaining quantitative data from interfaces suitable for such a comparison places the most severe demands upon high resolution imaging, when precision and accuracy are sought. Delocalisation can be overcome to some extent by aberration correction and/or exit wave reconstruction. But while the theory of exit wave reconstruction is well established, its application to interfaces in thicker foils is difficult. This paper considers these questions through a series of investigations using aberration corrected TEM, HAADF imaging and through focal series reconstruction of interfaces in Si3N4 and SrTiO3, in which the atomic structure of interfaces is sought to correlate with properties. In SrTiO3, long flat {100} GBs are frequently observed (e.g. Figure 1), while general GBs in niobium-doped samples are found to have a high density of {110} and {100} segments, their length depending upon sintering time (Figure 2) [1]. The termination of these faces is not yet clear, but it is likely to play a major role in facilitating grain growth, including abnormal grain growth. Combined HAADF and HRTEM imaging has been carried out to determine the terminations of these segments, and the atomic positions at the boundaries. In bicrystals, high resolution images have allowed the atomic structure, including any relaxation, to be studied and compared with alternative atomic model structures. In Si3N4, dopant atoms were shown to segregate to specific crystallographic sites at the IGF/GB interface [2]. However the average three dimensional coordinates of these sites determined from HAADF imaging are not so far in agreement with the positions determined by atomistic modelling. 1 2 3
S. J. Shih, K. J. Dudeck, S. Y. Choi, M. Baeurer, M. Hoffmann and D. J. H. Cockayne, Journal of Physics: Conference Series 94, (2008) 012008 G.B. Winkelman, C. Dwyer, T.S. Hudson, D. Nguyen-Manh, M. Döblinger, M and D.J.H. Cockayne, Philosophical Magazine Letters, 84/12 (2004) 755 Financial support from the European Commission under contract Nr. NMP3-CT-2005013862 (INCEMS) is acknowledged.
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b)
<110>
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G3
<100> G2 G1
500 nm
5 nm
Figure 1. (a) Typical low magnification image of a straight GB in undoped SrTiO3 sintered for 1 h at 1350 °C in air. (b) and (c) HREM images of the boundaries arrowed. Scale bars 500 nm and 5 nm.
Figure 2. Interface roughness of the {110} and {100} grain boundaries in Nb-doped SrTiO3 samples as a function of sintering time. Samples sintered at 1420°C in oxygen. [2]
Figure 3. High-resolution aberration-corrected image of pure STO with Sr/Ti~0.996, showing planar {100} GB interface. Scale bar = 200 nm
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Structure and chemistry of nanometer-thick intergranular films at metal-ceramic interfaces W.D. Kaplan and M. Baram Department of Materials Engineering, Technion – Israel Institute of Technology, Haifa, 32000 Israel [email protected] Keywords: Order, diffuse interfaces, HRTEM
The formation and existence of intergranular films (IGFs) have been explained over the years by different theories [1-7]. Avishai et al. showed a positive and relatively large Hamaker coefficient for amorphous films based on Si and Ca at Cu-Al2O3 and Ni-Al2O3 interfaces [8-9], which corroborates a study by Scheu et al. [10]. Due to questions regarding oxidation of the metal, this study verifies the existence of IGFS at Au-Al2O3 interfaces. The films were formed using a novel experimental approach, in which thin sputtered films of Au were dewetted on a sapphire substrate which was previously partially wetted with drops of anorthite glass (CaO-2SiO2-Al2O3) [11] (see figure 1). The process resulted in equilibrated metal particles residing on glass drops and on the sapphire substrate adjacent to the glass drops. The metal particles which formed on the glass drops ‘sank’ through the glass and reached the interface with the substrates. Aberration corrected transmission electron microscopy and analytical transmission electron microscopy were used to confirm the existence of the ~1nm thick amorphous films at the metal particle substrate interfaces (see figure 2). In addition, positive and relatively large Hamaker constants were calculated for the Au-film- Al2O3 interface, which indicates the existence of a stabilizing attractive van der Waals force, similar to equilibrium films at grain boundaries in ceramics. Adjacent to the glass drops, a ~1nm thick surficial film was also detected on the (0001) surface of sapphire substrates partially wetted by anorthite glass. 1. 2. 3.
R. Pandit, M. Schick and M. Wortis, Phys. Rev. B 26 (1982) 5112. J.W. Cahn, J. Chem. Phys. 66 (1977) 3667. R. Cannon, M. Hoffmann, R.H. French, H. Gu, A.P. Tomsia, E. Saiz, in "Grain Boundary Engineering in Ceramics", Ed. T. Sakuma, L.M. Sheppard, Y. Ikuhara, (2000) 427. 4. D.R. Clarke, J. Am. Ceram. Soc. 70 (1987) 15. 5. D.R. Clarke, T.M. Shaw, A.P. Philipse and R.G. Horn, J. Am. Ceram. Soc. 76 (1993) 1201. 6. S.J. Dillon, M. Tang, W.C. Carter and M.P. Harmer, Acta Mat., 55 (2007) 6208. 7. M. Tang, W.C. Carter and R.M. Cannon, Phys. Rev. B, 73 (2006) 024102. 8. A. Avishai, C. Scheu and W. D. Kaplan, Acta Mat., 53 (2005) 1559. 9. A. Avishai and W.D. Kaplan, Acta Mat., 53 (2005) 1571. 10. C. Scheu, G. Dehm and W.D. Kaplan, J. Am. Ceram. Soc., 84 (2001) 623. 11. M. Baram and W.D. Kaplan, J. Mater. Sci., 41 (2006) 7775.
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Figure 1. Secondary electron SEM of equilibrated gold particles in contact with the basal plane of sapphire, previously (partially) wetted with anorthite glass droplets. The arrowed particle was prepared for TEM (figure 2) using in-situ lift-out in a dual-beam FIB.
Figure 2. Aberration corrected (Cs=2.8μm) HRTEM micrograph of the interface between sapphire and a gold particle, containing a 1.1nm thick IGF.
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Studying nanocrystallization behaviour of different inorganic glasses using Transmission Electron Microscopy Somnath Bhattacharyya1, Th. Höche1, C. Bocker2, C. Rüssel2, A. Duran3, N. Hémono3, F. Muñoz3, M.J. Pascual3, K. Hahn4, P.A. van Aken4 1. Leibniz-Institut für Oberflächenmodifizierung e.V., D-04103 Leipzig, Germany 2. Otto-Schott-Institut für Glaschemie, Friedrich-Schiller-Universität Jena, Fraunhoferstraße 6, D-07743 Jena, Germany 3. Instituto de Cerámica y Vidrio (CSIC), Kelsen 5, 28049 Madrid, Spain 4. Max-Planck-Institut für Metallforschung, Heisenbergstr. 3, D-70569 Stuttgart, Germany [email protected] Keywords: TEM, nanocrystallization, phase separation
Nano technology and nano materials are considered to be the key technologies of the 21st century. Among the nano materials, glass ceramics are expected to play a major role, especially for optical applications. Glass-ceramics containing metal fluorides crystals with crystalline sizes 5 to 50 nm are considered to be materials with high potential for numerous photonic applications [1,2,3]. Crystals in this size range and narrow size distribution can only be obtained in multicomponent systems when the interface formed during nucleation and crystal growth acts as diffusional barrier that restricts crystal growth. In this present work, using transmission electron microscopy (TEM) we studied the nanocrystallization behaviour of two glass systems of silicate compositions in one of which (1) crystalline barium fluoride and in other (2) lanthanum fluoride precipitated. The two glass systems studied showed entirely different nanocrystallization behaviour. The glass with the nominal composition 40 SiO2 – 30 Al2O3 – 18 Na2O – 12 LaF3 turned out to possess lanthanum-rich phase-separation droplets that could not only be seen in the La distribution (Figure 1a), but – due to the very high scattering crosssections of La – also give rise to clearly discernable droplets in bright-field imaging (Figure 1b). During annealing at 645°C for 20 h these phase separation droplets were transformed into nanocrystals. The 7.52 Al2O3 – 15.04 K2O – 1.88 Na2O – 69.56 SiO2 – 6 BaF2 glass does not posses any phase separation while the glass ceramics (crystallisation at 600°C for 20 h) shows BaF2 nanocrystals of narrow size distribution as shown in Figure 2a. The Si mapping of this glass ceramics reveal that the interface of the crystalline phase is enriched with silica as shown in Figure 2b though due to poor spatial resolution (~3 nm) the inner crystalline BaF2 and the peripheral interface can not be separated. From the glass-chemistry point of view, silicon enrichment might well be responsible for the growth restrictions since it is expected to increase viscosity and hamper diffusion. This result is the direct experimental evidence of the nanocrystallization behaviour of glass ceramics proposed by Rüssel [3].
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S. Tanabe, H.Hayashi, T. Hanada and N. Onodera, Opt. Mater. 19 (2002), p. 343. Y. Wang and J. Ohwaki, Appl. Phys. Lett. 63 (1993), p. 3268. C. Rüssel, Chem. Mater. 17 (2005), p. 5843. Funding by the European Commission under contract Nr. NMP3-CT-2006-033200 is acknowledged.
(a)
(b)
Figure 1. (a) La-map of the phase-separated 40 SiO2–30 Al2O3–18 Na2O–12 LaF3 glass obtained by 120 kV energy-filtered TEM. (b) Zero loss filtered (using 15 eV energy slit around the Zero Loss peak) bright field image.
(a)
(b)
Figure 2. (a) Zero loss filtered bright field image of 7.52 Al2O3 – 15.04 K2O – 1.88 Na2O – 69.56 SiO2 – 6 BaF2 glass (b) Si map of the same area using 120 kV energyfiltered TEM.
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HRTEM and Diffraction Analysis of Surface Phases in Nanostructured LiMn1.5Ni0.5O4 Spinel F. Cosandey1, N. Marandian Hagh2 and G.G. Amatucci2 1. Department of Materials Science and Engineering, 2. Energy Storage Research Group, Department of Materials Science and Engineering, Rutgers University, Piscataway NJ 08854 [email protected] Keywords: HRTEM, Diffraction, LiMn1.5Ni0.5O4,
The Ni doped transition metal spinel LiMn1.5Ni0.5O4 is an excellent candidate as Liion cathode material for high voltage applications due to its high capacity (> 130 mAh/g) with good cycling and rate capabilities at room temperature. In order to improve its performance at high temperatures (60 oC), we have introduced a surface modification process involving treatments in hydrofluoric (H-treated) or phosphoric (Ptreated) acids followed by high temperature annealing. The surface modified LiMn1.5Ni0.5O4 spinel show drastically improved impedance and capacity retention with enhanced cycle life and rate capabilities over untreated spinel. In this paper, we present HRTEM, diffraction and image simulation results on the formation of a surface phase induced by these surface acidic and annealing treatments. The detailed description of sample preparation methods are presented elsewhere [1]. Electron diffraction (ED) and high resolution images (HRTEM) were obtained using a JEOL JEM 2010F operating at 200 kV. Kinematical diffraction and dynamical HRTEM images simulations were performed using JEMS program [2]. An electron diffraction pattern taken from a standard (STD) LiMn1.5Ni0.5O4 nanoparticle (50 nm) along the [110] zone axis is shown in Figure 1a with HRTEM image and corresponding FFT pattern shown in Figure 1b and 1c respectively. The results are consistent with a spinel Fd-3m structure as shown from the simulated diffraction pattern of Figure 3a. The electron diffraction pattern taken from a P-treated LiMn1.5Ni0.5O4 nanoparticle (50 nm) taken along the [011] zone axis is shown in Figure 2a with HRTEM image and corresponding FFT pattern shown in Figure 2b and 2c respectively. Extra superlattice reflections with (110), (210), (310) type reflections are observed in this P-treated sample. Similar results are observed for the H-treated sample. Kinematical electron diffraction simulations from the ordered P4332 spinel [3] phase shown in Figure 3c as well as HRTEM dynamical simulations are consistent with observed experimental results. 1 2 3
N. Marandian Hagh and G.G Amatucci, ECS Transactions, 11, (2007) http://cimewww.epfl.ch/people/stadelmann/jemsWebSite/jems.html. J.-Kim et al., Chem. Mater. 16, (2004), p.906.
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Figure 1 (a) [110] zone axis diffraction pattern, (b) HRTEM and (c) corresponding FFT of STD LiNi0.5Mn1.5O4.
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b
c
Figure 2. (a) [110] zone axis diffraction pattern, (b) HRTEM and (c) corresponding FFT of P-treated LiMn1.5Ni0.5O4 showing extra superlattice reflections.
a
b
Figure 3. Simulated nanoprobe [110] zone axis patterns for spinel LiMn1.5Ni0.5O4 with (a) Fd-3m and (b) P4332 space groups
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The structural origin of the antiferroelectric properties and relaxor behavior of Na0.5Bi0.5TiO3 V. Dorcet2, G. Trolliard1 and P. Boullay2 1. Laboratoire Sciences des Procédés Céramiques et des Traitements de Surface, SPCTS - CNRS-UMR-6638, 123 Av. Albert Thomas, F-87060 - Limoges Cedex (France). 2. Laboratoire de Cristallographie et Sciences des Matériaux, CRISMAT - CNRS UMR 6508, 6 Bd. du Maréchal Juin, F-14050 - Caen Cedex (France). [email protected] Keywords: TEM, NBT, phase transitions, ferroelectrics
Na0.5Bi0.5TiO3 (NBT) has been thoroughly investigated during the past forty years, mainly for its interesting ferroelectric and dielectric properties. At ambient temperature, this ferroelectric perovskite compound is rhombohedral with the polar R3c space group. It is known to transform to a tetragonal phase over a large domain of temperature of about 150°C during which the two phases coexist, which makes NBT an unusual compound exhibiting very specific behavior. In addition, if we compare the temperatures of phase transition deduced on one hand from the structural studies and on the others from the numerous experimental physical properties obtained over the last forty years, no clear correspondence exists. Then, between 200°C and 320°C, the correlation between the properties and the structure is not yet understood and a temperature in-situ analysis was carried out by transmission electron microscopy (TEM) to reconsider this phase transition. This study shows for the first time that the rhombohedral to tetragonal phase transition in NBT is in fact a two steps phase transitions. The transformation begins by a first order phase transition involving the reconstructive transformation of the rhombohedral (a-a-a- tilt system) phase into an orthorhombic one, through the formation of an intermediate modulated phase. This phase transition begins slightly over 200°C by the disappearance of the ferroelectric-ferroelastic domains (Figure 1). The intermediate modulated phase is observed up to 300°C, temperature at which it disappears (Figure 1). It is formed of Pnma orthorhombic sheets, appearing within the R3c matrix and exhibiting a-a+a- octahedra tilting system. These sheets are twin boundaries between two R3c ferroelectric domains. As the temperature increases, a micro-twining process takes place, generating more and more Pnma sheets. The increasing amount of Pnma sheets, in which the cations are displaced along [u0w]p, leads to a re-orientation of the polar vector within the R3c perovskite blocks initially orientated along [111]p. Thus, due to the pseudo merohedral twining law, two successive ferroelastic-ferroelectric R3c domains present polar vectors orientated in opposite direction, in a plane perpendicular to the modulated direction. The modulated phase explains the antiferroelectric property of NBT in this temperature range. The modulated phase is also at the origin of the relaxor behavior of NBT. In the close neighborhood of the Pnma sheets, the cations may occupy two kinds of atomic positions which are nearly equivalent from the point of view of their energy. These two S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 527–528, DOI: 10.1007/978-3-540-85226-1_264, © Springer-Verlag Berlin Heidelberg 2008
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positions are defined by the local structure of the R3c blocks ([111]p displacement) and the Pnma sheets ([u0w]p displacement). The cations could then alternatively jump from one position to the other by a flip-flop mechanism, giving rise to the so-called relaxor behavior of NBT. This study reconciles the structural properties together with the electrical behavior, showing that NBT is not so peculiar compound as it seemed.
Figure 1. Microstructural evolution of NBT viewed along the [001]p zone axis as a function of temperature. a) 20°C, b) 150°C, c) 200°C, d) 230°C, e) 241°C, f) 264°C, g) 274°C et h) 300°C.
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Electron beam probing of insulators H.-J. Fitting1, N. Cornet2, M. Touzin3, D. Goeuriot2, C. Guerret-Piécourt4, D. Juvé4, and D. Tréheux4 1. Physics Dept., Rostock University, Universitaetsplatz 3, D-18051 Rostock, Germany 2. Ecole Nationale Supér. des Mines, 158 cours Fauriel, F - 42023 Saint-Etienne, France 3. Université des Sciences et Technologies de Lille, F-59655 Villeneuve d'Ascq, France 4. Ecole Centrale de Lyon, 36 avenue Guy de Collongue, F - 69134 Ecully, France [email protected] Keywords: Non-conductive samples, Charge injection, Silica, Alumina
Electron beam irradiation and charge injection associated by selfconsistent charge transport in insulating samples are described by means of an electron-hole flight-drift model (FDM) implemented by an iterative computer simulation [1,2]. Ballistic scattering and transport of secondary electrons and holes is followed by electron and hole drift, their possible recombination and/or trapping in shallow and deep traps. Furthermore a detrapping by the temperature- and field-dependent Poole-Frenkel-effect becomes possible allowing even a charge hopping transport. In this context a special surface layer has been installed to investigate surface leakage currents, see Fig.1. As a main result the spatial distributions of currents j(x,t), charges ρ(x,t), electric field F(x,t), and potential V(x,t) are obtained in a selfconsistent procedure as well as the time dependent secondary electron emission rate σ(t) and surface potential V0(t) both experimentally accessible. For bulk full insulating samples the above quoted timedependent distributions approach the final stationary state under the condition j(x,t) = const = 0 and σ = 1. In case of remarkable surface leakage current iS the steady stationary final state is obtained with σ < 1, see Fig.2. The difference is collected with the surface current iS and can be measured by a respective ring electrode, Fig.3. In spite of a negative surface potential V0 the charge beneath the surface is positive due to the favored SE escape. But, generally we obtain a plus-minus-plus-minus spatial charge distribution with prevailing minus parts within the bulk insulator produced by a bipolar field: a positive field near the surface and a negative one in the remaining bulk. Thus, due to drift processes we obtain two opposite charge separations leading to the given quadro-polarized charge structure across the sample depth. The determine stationary-instationary sample current iSC(t) measurements allows us to measure the instationary charge displacement (incorporation) and polarization current ip as well as the surface leakage current iS, Figs.3 and 4. It offers the opportunity to investigate and to characterize insulating materials with respect to their quality, their electrical behavior, and radiation resistance. 1. 2.
M. Touzin, D. Goeuriot, C. Guerret-Piécourt, D. Juvé, D. Tréheux, H.-J. Fitting, J. Appl. Phys. 99 (2006) 114110-1-14 N. Cornet, D. Goeuriot, C. Guerret-Piécourt, D. Juvé, D. Tréheux, M. Touzin, H.-J. Fitting, J. Appl. Phys.103 (2008) in print
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Figure 1. Primary electron (PE) injection into a layered target; iBE : backscattered electron, iSE : secondary electrons, iS : surface current, ic : possible real conduction current, ip : polarization displacement current.
Figure 2. Total SE rate σ(t) and surface potential V0(t) as a function of irradiation time t and sample temperature T.
Figure 3. Sample shielding and measurements of sample current iSC, surface current iS, and polarisation current ip.
Figure 4. Simulated and experimental sample currents iSC (t) as given in Figure 3.
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Characterization of Ge-based clathrates oxidized in air by means of TEM and SEM C. Hébert1, B. Bartova1, M. Cantoni2, U. Aydemir3 and M. Baitinger3 1. LSME, Ecole polytechnique Fédérale, Station 12, CH-1015 Lausanne, Switzerland 2. CIME, Ecole polytechnique Fédérale, Station 12, CH-1015 Lausanne, Switzerland 3. Max Planck Instiute for Chemical Physics of Solids, Nöthnitzer Str. 40, D-01187 Dresden, Germany [email protected] Keywords: clathrates, oxide layer (or oxidation), TEM
Germanium based clathrate-I compounds are considered for their potential applications due to possible thermoelectric and optoelectronic properties. Clathrate compounds may be an alternative to well known thermoelectric materials at higher temperature. In this contribution we examined the stability of these compounds annealed in air. Such experiments are important to simulate the compounds behavior under the conditions of a technological process, e.g. in a thermoelectric module. Ba8Ni4Ge42 compounds were studied after the oxidation reaction. The compounds considered here have a well-defined composition before oxidation and show no superstructure. The clathrate has been crushed into a powder and sieved to 63 micrometers. The oxidation reactions were performed in air for two different temperatures (645 K and 773 K) and for different times, monitored with X-ray diffraction (XRD). The oxidized samples have been investigated by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). For SEM observations the samples were embedded in resin and cut with an ultramicrotome. The resulting clean surface was imaged. The samples oxidized at 645 K have been investigated by XRD after 6h, 12h, 24h, 48h, 72h and 96h. Decomposition was not observed until 96h. However, an oxide layer covering the surface was detected in SEM for the samples annealed for 24 h and longer. The oxide layer for shorter times of annealing was not very well detectable. The thickness of the layer increases with the oxidation time. The mean values presented in Fig.1a are averaging over 10 measurements. Typical microstructure of the oxide layer is shown in Fig.1b. From one grain of the sample annealed for 72h TEM lamella was prepared using focused ion beam. It shows a 100-150 nm thick amorphous layer of Ge-oxide with small crystalline inclusions of Ni-rich phases (most probably Ni oxide) at the interface with the clathrate (Figure 2). The external part layer of the amorphous phase has a composition close to GeO2, while the part at the interface with the clathrate is close to GeO. The Ba concentration is very homogenous in the amorphous phase. The Ba/Ge atomic ratio is similar to the one of the original phase suggesting very little or no Ba diffusion. The oxidation process at 773 K is much faster than at 645 K. The structure of the oxide layer is also quite different and much more porous as can be seen from figure 3. S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 531–532, DOI: 10.1007/978-3-540-85226-1_266, © Springer-Verlag Berlin Heidelberg 2008
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The increase in oxidation rate with time suggests that the Ni segregation plays a key role in the oxidation process. This is confirmed by the fact that the clathrates with lower Ni concentration are more stable in air.
Figure 1. (a) Thickness of the amorphous layer as a function of annealing time for the treatment at 645 K. (b) Typical SEM image of the amorphous layer in SEM, sample annealed for 72 h.
Figure 2. (a) DF-STEM image of the FIB lamella. (b) EDX chemical map of the squared region showing the oxidized layer and the Ni segregation at the interface between the amorphous layer and the clathrate.
Figure 3. SEM image of the sample oxidized 1h at 773 K showing the porous layer on the surface of the clathrate.
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Microstructure analysis of thin Cr2AlC films deposited at low temperature by magnetron sputtering Riza Iskandar1, Darwin P. Sigumonrong2, Jochen M. Schneider2 and Joachim Mayer1 1. Central Facility for Electron Microscopy, RWTH-Aachen University, Ahornstr. 55, 52074 Aachen, Germany 2. Materials Chemistry, RWTH Aachen University, Kopernikusstr. 16, 52074 Aachen, Germany [email protected] Keywords: Cr2AlC; Transmission electron microscopy
Cr2AlC belongs to the so called of MAX phases which have promising features because they combine unique properties usually found in either metals or ceramics. They have high elastic moduli and they are machinable [1], they exhibit good damage tolerance [2,3] good corrosion and oxidation resistance[2,4], good electrical and thermal conductivity[1], they passes a low friction coefficient, good damage tolerance [2] and high temperature stability [3]. In this work, Cr2AlC films have been synthesized by magnetron sputtering from a compound Cr2AlC target [5,6,7]. In this process, the influences of different substrate bias, deposition temperature (Ts), deposition time (t) and substrate material on the structure evolution have been investigated. The microstructures of the films have been studied using transmission electron microscopy. For this investigation, high-resolution transmission electron microscopy (HRTEM) and also energy filtering transmission electron microscopy (EFTEM) have been applied. Our investigations show that only the variation of deposition temperature and time influence the crystallization process of the films. A very thin amorphous interlayer (less then 1 nm) and comparatively large grains sizes have been observed in the samples grown at 450°C and 650°C. The microstructure of sample grown at 450°C is columnar with an average grain size of 80 nm. The sample grown at 650°C, besides having columnar grains too, has bigger grain size (130 nm) with many crystal orientations. High-resolutions images show that the films contain defects mainly twin boundaries and stacking faults. EFTEM images indicate that elemental distributions in the films are homogenous and no significant incorporation of oxygen has occurred [8]. 1. 2. 3. 4. 5.
M.W. Barsoum, T. El-Raghy, J. Am. Ceram. Soc. 79 (1996) 1953. M.W. Barsoum, T. El-Raghy, L. Ogbuji, J. electrochem. Soc. 144 (1997) 2508. T. El-Raghy, M.W. Barsoum, A. Zavaliangos, S. Kalidindi, J. Am. Ceram. Soc. 82 (1999) 2855. Z. Sun, Y. Zhou, S. Li, Acta Mater. 49 (2001) 4347. O. Wilhemlmsson, J.-P. Palmquist, T. Nyberg, U. Jansson, Appl. Phys. Lett. 85 (6) (2004) 1006
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 533–534, DOI: 10.1007/978-3-540-85226-1_267, © Springer-Verlag Berlin Heidelberg 2008
534 6. C. Walter, D.P. Sigumonrong, T. El-Raghy, J.M. Schneider, Thin Solid Films 515 (2006) 389393 7. J.M. Schneider, D.P. Sigumonrong, D. Music, C. Walter, J. Emmerlich, R. Iskandar and J. Mayer, Scripta Mate. 57 (2007) 1137-1140. 8. Support of these investigations by the DAAD and the DFG is gratefully acknowledged.
a
b
c
d
Figure 1. HRTEM images showing twin boundaries and stacking faults of specimen grown at 450°C (a,b) and 650°C (c,d).
BF
Cr
Al
C
O
Figure 2. EFTEM elemental maps showing elemental distributions of the sample. The films are homogenous and no significant incorporation of oxygen can be detected.
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Study of structural variation in YBaCo4O7+δ by electron diffraction Y. Jia1, H. Jiang1, M. Valkeapää2, M. Karppinen2 and E.I. Kauppinen1 1. Laboratory of Physics and Center for New Materials, 2. Laboratory of Inorganic Chemistry, Department of Chemistry, Helsinki University of Technology, FIN-02150 Espoo, Finland [email protected] Keywords: electron diffraction, modulated structure, YBaCo4O7+δ
In recent years, the layered cobalt oxide, YBaCo4O7+δ, has attracted much attention due to its interesting structural, electronic and magnetic properties. The basic structure of the oxygen-stoichiometric YBaCo4O7 has been described as a mixture of hexagonal closed packing (hcp) and cubic closed packing (ccp) with oxygen atoms stacking along c axis in the sequence –A-B-C-B- [1], which consists of alternating layers of Kagomé and triangular lattices of cobalt atoms [2]. The Kagomé and triangular layers are made up of the corner-sharing oxygen tetrahedra with the centres occupied by Co2+/Co3+ ions in the ratio of 3/1. In this work, the structure of YBaCo4O7+δ prepared by the solid state reaction was studied by means of electron diffraction. Compared with its parent phase YBaCo4O7, a phenomenon of structural variation due to introducing of extra oxygen in the structure has been observed. Origin of structural modulation is discussed. Figure 1 shows a tilt-series (around c axis) electron diffraction patterns (EDPs) of the parent phase YBaCo4O7 along zone axes of (a) [110], (b) [230], (c) [120], (d) [130], (e) [140] and (f) [010], and Figure 2 presents EDPs of an oxygenated compound YBaCo4O7+δ along zone axes of (a) [032], (b) [031], (c) [3-32], (d) [1-11], respectively. The electron diffraction patterns have been indexed as a hexagonal structure with a unit cell of ah = 6.4 Å and ch =10.2 Å. Compared with Figure 1, it is interesting to remark that in Figure 2 there are systematic satellite reflections appearing around main reflections in all diffraction patterns, which indicates that the crystal structure of the oxygenated phase YBaCo4O7+δ is modulated. From their EDPs, the modulation vector is determined as q = (a* + b* ) / 3 . This is a commensurate modulation which can also be described as a hexagonal superstructure with a super unit cell of parameters a s = 3a h and c s = ch . Obviously, the structural variation in the oxygenated sample results from the introducing of extra oxygen into the parent phase YBaCo4O7. The extra oxygen atoms may form edge-sharing Co-O octahedra with Co3+ ions [3], which is considered to account for the origin of the structure modulation. 1. 2. 3. 4.
M. Valldor and M. Andersson, Solid State Sci. 4 (2002), p.923. M. Soda, Y. Yasui, T. Moyoshi, et.al., J. Magn. Magn. Mater. 310 (2007), e441. M. Valldor, in "New Topics in Condensed Matter Research", ed. J. V. Chang, Nova Science Publishers, New York, (2007), p. 75-102. We acknowledge financial supports from CIMO and the Academy of Finland.
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Figure 1. c axis tilt-series electron diffraction patterns of the parent phase, YBaCo4O7, with zone axes of (a) [110], (b) [230], (c) [120], (d) [130], (e) [140] and (f) [010], respectively.
Figure 2. Electron diffraction patterns of YBaCo4O7+δ along the zone axes of (a) [032], (b) [031], (c) [3-32], and (d) [1-11], respectively.
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Exsolution phenomena in glass-ceramic systems I. Tsilika, Ph. Komninou, G.P. Dimitrakopulos, Th. Kehagias and Th. Karakostas Department of Physics, Aristotle University of Thessaloniki, 54124 Thessaloniki, Greece [email protected] Keywords: wollastonite, diopside, exsolution, glass-ceramics
Exsolution is a known mechanism of phase transformation in pyroxenes. In particular, the transformation of the pyroxenes augite to pigeonite and vice-versa has been a subject of extent and thorough study since 1941 [1]. The transformation process involves the formation of nuclei, which gradually grow in lamellae during cooling. The precise orientation of the lamellae in relation to the parent structure is determined by the minimum interfacial energy. Transmission electron microscopy studies in these systems reveal submicroscopic textures such as periodic exsolution and micro-twinning [2]. In the present study, the phase transformation of wollastonite to diopside via the mechanism of exsolution is reported, by means of conventional and high-resolution transmission electron microscopy (TEM-HRTEM) and X-ray energy dispersive spectroscopy (EDS) microanalysis. This is a pyroxenoid to pyroxene transformation that takes place via cation substitution, i.e. a diffusion controlled process. The exact mechanism of the pyroxenoid-pyroxene transformation, which is caused by changes in pressure, temperature and cation size, has not been elucidated so far [3]. The glass-ceramics examined belonged to the SiO2-Na2O-CaO system with an initial batch composition of 55 SiO2-5 Na2O-25 CaO (wt%) and were produced through the vitrification process [4]. The remaining 15% were smaller amounts of other oxides (ZnO, Fe2O3 etc.) that have been added as nucleates, in order to obtain volume crystallization. The powder mixture was initially homogenized and then casted at 1400oC. The vitreous into glass-ceramics transformation was obtained by a two stage thermal treatment, i.e. nucleation at 680oC for 15 min, followed by crystallization at 900oC for 2h. The electron diffraction analysis in TEM showed wollastonite as the major crystalline phase formed in its two most frequent forms, i.e. triclinic with a P 1 space group and monoclinic with a P21/m space group. Diopside (space group C2/c) was found to exsolve from the parent structure only in the areas of monoclinic wollastonite, where the concentration of Fe and Zn is considerably high, as it was detected from the EDS microanalysis. The elements of Fe, Zn are divalent cations that can be accommodated in the large octahedral sites of Ca in the wollastonite structure by ionic substitution. The substitution of the large cations of Ca with the smaller cations of Fe and Zn, in large amounts, influenced the periodicity of the silicate chain and induced the formation of the clinopyroxene structure. In the early stages, diopside formed as nanosized nuclei, which were apparent in HRTEM images by moiré fringes. The HRTEM image of a nanoparticle of diopside along with the corresponding electron diffraction pattern (EDP) is shown in Figure 1(a,b). Subsequently, diopside grew as lamellae inside the parent matrix of wollastonite illustrated in Figure 1(c). S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 537–538, DOI: 10.1007/978-3-540-85226-1_269, © Springer-Verlag Berlin Heidelberg 2008
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1. 2. 3. 4.
H. Hess, Amer. Miner. 67 (1941), p. 573 H. Ried and H. Fuess, Phys. Chem. Miner. 13 (1986), p. 113. D. Veblen, Amer. Miner. 70 (1985), p. 885. P. Kavouras, Th.A. Ioannidis, Th. Kehagias, I. Tsilika, K. Chrissafis, S. Kokkou, A. Zouboulis, Th. Karakostas, J. Eur. Ceram. Soc. 27 (2007), p. 2317.
(b)
(a)
(c) Figure 1. (a) Early stages of diopside formation; moiré fringes arise due to superposition of a diopside nanoparticle and monoclinic wollastonite structure; (b) corresponding EDP, where the [ 1 1 2] zone axis of diopside (intense spots) is superimposed to the [101] zone axis of the parent wollastonite. (c) Growth of diopside lamellae within the wollastonite parent matrix.
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Transmission Electron Microscopy Studies of Lead-Free Ferroelectrics in the System BNT-BT-KNN H.-J. Kleebe1, J. Kling2, L. Schmitt1, S. Lauterbach1, and H. Fuess2 1. Geomaterial Science, Institute for Applied Geosciences, Schnittspahnstrasse 9, 64287 Darmstadt, Germany 2. Structure Research, Materials Science Institute, Petersenstrasse 23, 64287 Darmstadt, Germany [email protected] Keywords: lead-free ferroelectrics, TEM, electron diffraction, superlattice reflections, cation disorder, octahedral tilting
Lead zirconate titanate (PZT) materials are widely used in industrial applications like fuel injectors or actuators. The possible environmental problems caused by the lead content of the conventional devices drives the need to replace lead-based systems by new, nontoxic components. However, those novel materials have to compete with the electrical and mechanical properties of to days PZT. Lead-free piezoelectric ceramics, (1−x−y)Bi0.5Na0.5TiO3–xBaTiO3–yK0.5Na0.5NbO3 (0.05≤x≤0.07 and 0.01≤y≤0.03) have been synthesized by a conventional solid state sintering method. The room temperature ferroelectric and piezoelectric properties of these ceramics were studied. In addition, microstructural features were investigated by transmission electron microscopy (TEM) in conjunction with electron diffraction. Based on the measured properties, the ceramics were categorized into two groups: group I compositions having dominant ferroelectric order and group II compositions displaying mixed ferroelectric and antiferroelectric properties at room temperature. A composition from group II near the boundary between these two composition groups exhibited a strain of approximately 0.45% at an electric field of 8 kV/mm. Polarization in this composition range was not stable because the piezoelectric coefficient d33 at zero electric field was as low as 30 pm/V. The converse piezoelectric response becomes weaker when the composition deviates from the boundary between the groups toward either the ferroelectric or antiferroelectric regime. The latter results were rationalized based on a field induced antiferroelectric-ferroelectric phase transition. Both x-ray and TEM investigations were not able to clearly reveal structural differences among the materials investigated, even when crossing the antiferroelectricferroelectric boundary. No prominent domain structure, as commonly seen in PZT polycrystals, was observed in either of the material groups. This observation suggested a cubic crystal structure, which was consistent with data obtained by X-ray diffraction. In contrast to the featureless grain structure found via TEM imaging, electron diffraction revealed weak superlattice reflections in prominent zone axis. The origin of these superlattice reflections, as also shown in the inset in Figure 1, can be twofold: (i) tilting of the oxygen octahedra (Glazer notation) or (ii) cation sublattice ordering. Details of the TEM analysis and electron diffraction study will be presented and discussed in light
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 539–540, DOI: 10.1007/978-3-540-85226-1_270, © Springer-Verlag Berlin Heidelberg 2008
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of high-resolution TEM imaging (and STEM analysis) of these rather complex BNTBT-KNN structures. A correlation between crystal structure and electro-mechanical response will be presented.
[001
90-06-04
Figure 1. TEM micrograph of a lead-free ferroelectric sample of composition 90-06-04. The inset of the electron diffraction pattern clearly reveals the presence of superlattice reflections. The arrows indicate the presence of dislocations, which are commonly present in these materials. 1. S.-T. Zhang, A.B. Kounga, E. Aulbach, T. Granzow, W. Jo, H.-J. Kleebe, J. Rödel, "LeadFree Piezoceramics with Giant Strain in the System Bi0.5Na0.5TiO3− BaTiO3−K0.5Na0.5NbO3. I. Structure and Room Temperature Properties", J. Appl. Phys., 103 (2008) 034107-1-8.
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ReO3-related aluminum tungsten oxides showing a novel type of crystallographic shear structure F. Krumeich1 and G.R. Patzke2 1. Laboratory of Inorganic Chemistry, ETH Zurich, 8093 Zurich, Switzerland 2. Institute of Inorganic Chemistry, University of Zürich, 8057 Zürich, Switzerland [email protected] Keywords: structure determination, HRTEM, HAADF-STEM
A series of new aluminum tungsten oxides was found and structurally characterized by means of electron microscopy [1]. Symmetry and metric of these new phases were determined by selected area electron diffraction (Figure 1a). While the short crystallographic axis (~ 0.375 nm) approximately corresponds to the diagonal of a WO6 octahedron, the perpendicular axis shows five distinct superstructure reflections between the main reflections. From tilting experiments, the length of the axis in the third direction was determined to be ~ 0.375 nm as well. The observed reflection condition h + l = 2n for h0l points to a centered cell (Figure 1a). These findings are in accordance with a body-centered tetragonal unit cell (I4/mmm (space group No. 139); a ≈ 0.375, c ≈ 3.95 nm). The presence of Al in the investigated crystallites was confirmed by EDX spectroscopy. A structural model for Al4W10O32 was derived from HRTEM and HAADF-STEM images (Figure 1c,d) which both show the W positions (dark or bright patches, respectively). The structure consists of slabs of [5 x ∞ x ∞] corner-sharing WO6 octahedra that can be regarded as cuttings of the ReO3 type. Adjacent slabs are shifted by [½½½] with respect to each other as required by the body-centered unit cell (Figure 1b). The slabs of WO6 octahedra are connected via edge-sharing to AlO6 octahedra situated in the intermediate space. Simulated HRTEM images agree well with the experimental ones (insets in Figure 1c) and thus support the proposed structural model. The connection between adjacent slabs of WO3 via planes of AlO6 octahedra has been observed here for the first time and indeed represents a novel type of crystallographic shear operation for ReO3-type structures [2,3]. This crystal structure is very flexible since the width of the WO3 bands can vary. This leads to the formation polytypes with the general formula Al4W2nO6n+2 (n = 4-7). Frequently, the polytypes are intergrown with each other, and most crystallites are disordered and contain stacking faults (Figure 2). 1. 2. 3. 4.
F. Krumeich, G.R. Patzke, Microsc. Microanal. 13, Suppl. 3 (2007) 366. B. G. Hyde, M. O’Keefe, Acta Crystallogr. A29 (1973) 243. B.-O. Marinder, Angew. Chem. Int. Ed. 25 (1986) 431. We thank the EMEZ (electron microscopy ETH Zurich) for measuring time. HRTEM and SAED were performed on a CM30 (Supertwin lens, LaB6 cathode, Vacc = 300 kv); HAADFSTEM on a Tecnai F30 (SuperTwin lens, FEG, Vacc = 300 kv).
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 541–542, DOI: 10.1007/978-3-540-85226-1_271, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. (a) Electron diffraction pattern along [010]. The inset shows an enlarged section with some indices of the tenfold superstructure given. (b) Structural model of Al4W10O32 in projection onto the ac plane. (c) HRTEM image of Al4W10O32 along [010]. Simulations calculated with the structural data of Al4W10O32 for thicknesses of 1.5 nm (top) and of 4.5 nm (bottom) are shown as insets and marked by white lines (length ≈ 3.95 nm). (d) HAADF-STEM image revealing the W position as bright patches.
Figure 2. HAADF-STEM image of a disordered crystal region. The width of the ReO3type bands is indicated by the number of W positions.
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Structural Characterisation by TEM of a New Homologous Series Bi2n+4MonO6(n+1); n=3,4,5 and 6 A.R. Landa-Canovas1, J. Hernández-Velasco1, E. Vila1, J. Galy2 and A. Castro1 1. Dep. Sólidos Iónicos, Instituto de Ciencia de Materiales de Madrid, CSIC, Sor Juana Inés de la Cruz 3, Cantoblanco, 28049 Madrid, Spain. 2. Centre d’Elaboration de Matériaux et d’Etudes Structurales, CNRS, 29 rue Jeanne Marvig,B.P. 94347, 31055 Toulouse Cedex 4, France [email protected] Keywords: Bismuth-molybdenum oxides, fluorite structure, ionic conductors
Bi2O3-MoO3 system presents many phases with a fluorite-type related structure with anion vacancies and are, therefore, good oxygen conductors. Usually, this compounds are synthesized by the traditional ceramic method. However, the high temperatures employed by this method avoid the isolation of low-temperature phases. We have used low temperature methods like the so called n-butylamine method [1]. The precursors were obtained by precipitating stoichiometric amounts of Bi3+ over MoO3 (10 different mBi2O3-MoO3 compositions with m=1.0-1.7) by adding drop by drop (2mL/min) a 1M n-butylamine solution at room temperature heated at increasing temperatures until final temperatures between 600-700C [1]. Powder X-ray diffraction allows to identify four pure phases for Bi10Mo3O24, Bi6Mo2O15, Bi14Mo5O36 y Bi8Mo3O21 nominal compositions. Selected area electron diffraction, see Fig. 1, has been used to obtain lattice parameters of all these phases. All the patterns can be interpreted as a basic fluorite-type structure which gives rise to the basic lattice reflections. Weaker reflections can be interpreted as satellite reflections produced by a structural modulation with q= (602)*. =1/13, 1/16, 1/19 and 1/22 for the four different phases. Fig. 2 schemes how can we generate the different reciprocal lattices from the basic fluorite-type structure by applying the q vector of each phase. The relationship between the different phases shows the existence of a homologous series with general formula Bi2n+4MonO6(n+1) where n = 3, 4, 5 and 6. The relationship of the unit cell with respect to the basic fluorite-type structure is represented in Fig. 3. In Fig. 4 we show a HREM micrograph of the simplest phase, i.e., Bi10Mo3O24. From this HREM image and using the metric relationship between the observed supercell and the basic fluorite one we have deduced a cationic model for this new phase whose image simulation, inserted in Fig. 4, matches very well with the experimental image. Similar cationic models have been developed for the other three phases with satisfactory matching. Powder neutron diffraction confirms the proposed model, locates the oxygen anions and refines the structure with very good results. 1.
E. Vila, A.R. Landa-Cánovas, J. Galy, J.E. Iglesias y A. Castro; J. Solid St. Chem. 180 , 661 (2007)
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 543–544, DOI: 10.1007/978-3-540-85226-1_272, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. SAED patterns of phases Bi10Mo3O24, Bi6Mo2O15, Bi14Mo5O36 y Bi8Mo3O21
Figure 2. Scheme for the SAED patterns of phases Bi2n+4MonO6(n+1)
Figure 3. HREM micrograph and image simulation inserted of Bi10Mo3O24 phase.
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Structural characterisation of a new rich iron layered oxide TlεSr25-εFe30O76-x C. Lepoittevin1, S. Malo2, S. Hebert2, M. Hervieu2 and G. Van Tendeloo1 1. EMAT laboratory, University of Antwerpen, Groenenborgerlaan 171, 2020 Antwerpen, Belgium. 2. CRISMAT-ENSICAEN laboratory, University of Caen, Bd du Maréchal Juin, 14050 Caen cedex, France. [email protected] Keywords: iron oxide, transmission electron microscopy, layered structures
Iron based oxides are the source of rich chemistry thanks to the ability of iron to adopt different coordination environments and various oxidation states, from 2+ to 4+. For example, the structurally rich Sr-Fe-O system is characterised by numerous families of layered structures, generated thanks to the great adaptability of perovskite layers to other structural units. In the rich iron part of this system, the oxide Sr4Fe6O13-δ has been the subject of numerous papers reporting its properties for applications in the field of oxygen-selective membranes [1,2,3,4]. Recent papers described its structure as a stacking of a single perovskite layer with a complex oxygen deficient rock salt-type layer (RS), in which the oxygen stoichiometry is related to the incommensurate nature of the modulated structure [5,6,7,8]. This phase is the m=1 member of a new potential perovskite
family of structures [Sr2Fe2O7-δ]RS[Sr2Fe2O6 ]m
. By substituting strontium with
bismuth in Sr4Fe6O13-δ, the m=2 member Bi4Sr14Fe24O56 has been stabilized and structurally characterised by transmission electron microscopy and single crystal X-ray diffraction [9,10]. With the aim of stabilizing other members of this structural family, substitutions with thallium have been carried out, leading to a new strontium ferrite TlεSr25-εFe30O76-x (ε ≈ 0.5) which structure has been characterised by transmission electron microscopy. The electron diffraction study evidenced an orthorhombic sub-cell with asub-cell ≈ b ≈ 5.4 Å, c ≈ 42 Å and a F lattice. Satellites reflections in commensurate positions are observed on the [010] ED pattern (Figure 1) and the resulting modulation G G vector q = α a * ,with α = 2/5, allow to define a five-fold supercell with a ≈ 5×asubcell ≈ 27 Å with Bbab as a possible space group. G The [010] HRTEM image (Figure 1) evidences a regular stacking along c of a quadruple perovskite layer and an oxygen deficient rock salt-type layer. In the latter, the G complex variations of contrast along a illustrate the commensurate modulation as a regular alternation of two bright sticks with three weaker dots. Thanks to the previous results obtained on the structural study of Sr4Fe6O13-δ[5,8], these "2:3" sequences are correlated with a double ribbon of edge-sharing bi-pyramids [FeO5] alternating with two ribbons of mono-caped tetrahedra [FeO4] surrounding one ribbon of distorted tetragonal pyramids [FeO5]. S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 545–546, DOI: 10.1007/978-3-540-85226-1_273, © Springer-Verlag Berlin Heidelberg 2008
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The new ferrite TlεSr25-εFe30O76-x (ε ≈0.5) represents the m = 4 member of the family perovskite and opens the route to numerous of layered structures [Sr2Fe2O7-δ]RS[Sr2Fe2O6 ]m new materials. 1. 2. 3. 4.
H. Fjellvag, B.C. Hauback, R. Bredesen, J. Mater. Chem.. 7 (1997), p. 2415. T. Armstrong, S. Guggila, A. Manthiram, Mater. Res. Bull. 34 (1999), p. 837. S. Guggila, T. Armstrong, A. Manthiram, J. Solid State Chem. 145 (1999), p. 260. A. Fossadl, L.T. Sagdahl, M.A. Einarsrud, K. Wiik, T. Grande, P.H. Laresen, F.W. Poulsen, Solid State Ionics 143 (2001), p. 367. 5. B. Mellenne, R. Retoux, C. Lepoittevin, M.Hervieu, B. Raveau, Chem. Mater. 16 (2004), p. 5006. 6. M.D. Rossell, A.M. Abakumov, G. Van Tendeloo, J.A. Pardo, J. Santiso, Chem. Mater. 16 (2004), p. 2578. 7. M.D. Rossell, A.M. Abakumov, G. Van Tendeloo, M.V. Lomakov, S. Ya. Istomin, E.V. Antipov, Chem. Mater. 17 (2005), p. 4717. 8. O. Pérez, B. Mellenne, R. Retoux, B. Raveau, M. Hervieu, Solid State Sciences 8 (2006), p. 431. 9. C. Lepoittevin, S. Malo, M. Hervieu, D. Grebille, B. Raveau, Chem. Mater. 16 (2004), p. 5731. 10. D. Grebille, C. Lepoittevin, S. Malo, O. Pérez, N. Nguyen, M. Hervieu, J. Solid State Chem. 179 (2006), p. 3849.
Figure 1. [010] HRTEM image and electron diffraction pattern of TlεSr25-εFe30O76-x.
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EBSD studies of stress concentrations in ferroelectrics I. MacLaren1, M.U. Farooq1, R. Villaurrutia1, T.L. Burnett2, T.P. Comyn2, A.J. Bell2, H. Kungl3, M.J. Hoffmann3 1. Department of Physics and Astronomy, University of Glasgow, Glasgow G12 8QQ, UK 2. Institute for Materials Research, University of Leeds, Leeds LS2 9JT, UK 3. IKM, Universität Karlsruhe, Haid-und-Neu-Str. 7, 76131 Karlsruhe, Germany [email protected] Keywords: EBSD, ferroelectrics, stress
The junctions between lamellar domain boundary structures in so-called “herringbone structures” are expected to contain stress concentrations [1,2]. These herringbone domain junctions have been studied using electron backscatter diffraction to reveal crystallographic relationships across the domain boundaries and this data also reveals information about the localized stress concentrations at the junctions. “Figure 1a” shows an EBSD orientation map from one such herringbone domain junction in a 0.5 BiFeO3 – 0.5 PbTiO3 single crystal (an example of a similar junction from the same sample is shown in “Figure 1b”). The individual domains can be clearly identified in this map with just occasional misindexing (random colours) or failures to index (black points). The misorientations between the “90°” domains to either side of the junction correspond well to expectations – ab corresponds to a misorientation of about 84.5° about <100> and cd corresponds to a misorientation of 83° about <100>, as may be seen by reference to Figures 1c and 1d. These correspond well to the expected tilt values calculated from published lattice parameters for this material [3,4]. The misorientations across the junction have also been carefully measured and correspond well to crystallographic models. “Figure 2a” shows an EBSD orientation map from a herringbone domain junction in a ceramic of PbZr0.425Ti0.575O3 doped with 1 mol. % La and 2 mol. % Sr. This map shows the domain boundaries clearly to either side of the boundary and the tilt angles of these can be seen in the boundary map of “Figure 2b”. Again, misorientations correspond well to the lattice parameters measured by X-ray diffraction. In this latter case, how the tilt varies across the junction of the domains ad and a’d’ could be measured in detail due to the high spatial resolution of the map and this information allowed the estimation of internal stresses at the boundary of 0.56 GPa. Such high internal stresses are clearly of significance to properties and fatigue behaviour of the ceramics. 1 2 3 4
G. Arlt and P. Sasko, J. Appl. Phys., 51 (1980) p. 4956. I. MacLaren, L.A. Schmitt, H. Fuess, H. Kungl, and M.J. Hoffmann, J. Appl. Phys., 97, (2005) 094102. S.A. Fedulov, Yu.N. Venevtsev, G.S. Zhdanov, E.G. Smazhevskaya, and I.S. Rez, Kristallografiya, 7, (1962) p. 77. T.L. Burnett, T.P. Comyn, E. Merson, and A.J. Bell, IEEE T. Ultrason. Ferr., (2008) in press.
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Figure 1. SEM imaging and EBSD maps from a BF-PT sample: a) EBSD orientation map (coloured according to the inverse pole figure); b) backscattered image of a herringbone domain junction in the same sample; c) boundary map (top right) showing boundaries coloured according to d) the misorientation histogram.
Figure 2. EBSD mapping of a domain junction in La-Sr-doped PZT: a) Orientation map coloured according to the inverse pole figure showing clear lamellar domains; b) Boundary map showing the “90°” domain boundaries.
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High-resolution pictures of nucleation growth triangle of 180° ferroelectric domain wall in a thin film of LiTaO3 obtained by Lorentz DPC-STEM Takao Matsumoto, Masanari Koguchi & Yoshio Takahashi Central Research Laboratory, Hitachi, Ltd., 1-280 Higashi-Koigakubo, Kokubunji-shi, Tokyo 185-8601, Japan [email protected] Keywords: ferroelectric, domain wall, Lorentz DPC-STEM
Motions of 180° domain walls in ferroelectric materials are critically important in state-of-the-art applications of ferroelectric materials[1-3]. So far, such motions have been described by a classical triangular-nucleation model proposed by Miller and Weinreich[4]. Recently, based on first-principle calculations, a square-shaped nucleation with a diffuse boundary was proposed[5]. From the experimental point of view, the mechanism of the motions is still unclear owing to a lack of a visualization technique with sufficient spatial resolution. Following these works, in our present study, Using differential-phase-contrast mode of Lorentz scanning transmission electron microscopy (Lorentz DPC-STEM[6-8]), we obtained high-resolution images of a nucleation growth triangle of a charged 180° ferroelectric domain wall coupled head-toEvidence of charge accumulation head in a thin film of LiTaO3 single crystal. accompanying large strain due to the head-to-head coupling of polarization at the boundary was obtained. Moreover, 180° domain walls were found to be liable to move, so they often changed their appearance during observation on a time-scale of several tens of seconds. The wedge angle at the tip of nucleation-growth triangle, as described by the Miller–Weinreich model, was determined as 40°±2°. This angle plausibly corresponds to atomistic one-to-one sequences of alternating charged domain walls and uncharged anti-parallel domain walls. 1. 2. 3. 4. 5. 6. 7. 8. 9.
J. F. Scott, and C. A. P. d., Araujo, Science 246 (1989), p. 1400. J. F.Scott, J. Phys.: Condens. Matter 18 (2006), p. R361. N. Setter et al., J. Appl. Phys. 100 (2006), p. 051606. R. C. Miller and G. Weinreich, Phys. Rev. 117 (1960), p. 1460. Y.-H. Shin, I. Grinberg, I.-W.Chen, and A. M. Rappe, Nature 449 (2007), p. 881. N. H. Dekkers & H. de Lang, Optik 41 (1974), p. 452. J. N. Chapman, S. McVitie & S. J. Hefferman, J. Appl. Phys. 69 (1991), p. 6078. Y. Yajima et al., J. Appl. Phys. 73 (1993), p. 5811. We acknowledge the help of the Dr. Yusuke Yajima of Ibaraki University for his valuable advice and comments on Lorentz DPC-STEM. We also acknowledge the contribution of Dr. Nobuyuki Osakabe of Advanced Research Laboratory, Hitachi, Ltd., and Drs. Keigo Suzuki, Yasuhiro Motoyoshi, and Nobuyuki Wada of Murata Manufacturing Co., Ltd. for their numerous discussions throughout the present work.
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Figure 1. Reconstruction of polarization vector and charge density at the tip of nucleation triangle. a, The reconstructed polarization vector map. Directions polarization are designated as white arrows. A square-shaped tip is indicated by a red arrow-head. b, The reconstructed charge density. Faint lines of bright contrast at the wall indicate the localized positive charge. c, Definition of wedge angle. Bars represent 50 nm.
Figure 2. Schematic diagrams of the triangular spike-shaped 180° domain walls. a, The microscopic picture. b, the atomistic picture at the tip of the triangular nucleation wedge designated by the broken white circle in a. The angle of the wedge is determined by the alternating 180˚ domain wall coupled head-to-head carrying positive charge and parallel uncharged wall. Dimensions of the unit cell determines the wedge angle.
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Size and structure of barium halide nano-crystals in optically active fluorozirconate-based glasses P.T. Miclea1,2, B. Ahrens3, C. Eisenschmidt4, and S. Schweizer1,2 1. Fraunhofer Center for Silicon Photovoltaics, Walter-Hülse-Str. 1, D-06120 Halle, Germany 2. Institute of Physics, Martin-Luther-University of Halle-Wittenberg, HeinrichDamerow-Str. 4, D-06120 Halle, Germany 3. Department of Physics, Faculty of Science, University of Paderborn, Warburger Str. 100, D-33098 Paderborn, Germany 4. Institute of Physics, Martin-Luther-University of Halle-Wittenberg, Hoher Weg 8, D-06120 Halle, Germany [email protected] Keywords: barium chloride nano-crystals, fluorozirconate glass ceramic
Optically active fluorozirconate-based (FZ) glass ceramics offer a broad range of applications. The functionality of the glass ceramic can be modified by appropriate doping and thermal processing performed after the glass production. For use in digital radiography, for example, the FZ glass has been doped with europium and chlorine ions [1, 2]. Thermal processing by annealing in the vicinity of the glass transition temperature produces barium chloride nano-crystals in the glass. The glass ceramic can act either as a scintillator (able to convert ionizing radiation to visible light) [1], or as a storage phosphor (able to convert the radiation into stable electron-hole pairs, which can be read out afterwards with a scanning laser beam in a so-called “photostimulated luminescence” process) [2]. The scintillation and the photostimulated luminescence in these glass ceramics is caused mainly by the emission of divalent europium incorporated in hexagonal and orthorhombic barium bromide/chloride nano-crystals; divalent europium does not luminesce in the FZ glass matrix. Interestingly, the hexagonal phase is always formed first before it is converted into the orthorhombic phase. The hexagonal phase is used for scintillation, the orthorhombic one for storage phosphor applications. There is a clear tendency for larger particles at longer annealing times. This leads to a trade-off between optical transparency and the scintillation or storage efficiency. A combination of good optical transparency and efficiency depends on the annealing conditions used to induce crystallization. We investigated the particle size distribution of the barium chloride nano-crystals versus annealing temperature by means of transmission electron microscopy (TEM), the structure and morphology of the nano-crystals by high-resolution TEM, and the composition by energy dispersive x-ray (EDX) spectroscopy.
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1. 2. 3.
J.A. Johnson, S. Schweizer, B. Henke, G. Chen, J. Woodford, P.J. Newman, and D.R. MacFarlane, Journal of Applied Physics 100 (2006), 034701 J.A. Johnson, S. Schweizer, and A.R. Lubinsky, Journal of The American Ceramic Society 90(3) (2007), 693 S. Schweizer, L.W. Hobbs, M. Secu, J.-M. Spaeth, A. Edgar, G.V.M. Williams, Appl. Phys. Lett. 83(3) (2003), 449
Figure 1. TEM images of an Eu-doped fluorochlorozirconate glass-ceramic annealed at 290°C for 5 min (left) and 10 min (right) [3]. The scale marker is 100 nm in length.
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Domain Structure And Microstructure Development of BaTiO3 Doped With Rare-Earth Dopants V. Mitic1,2, V.B. Pavlovic3, V. Paunovic1, M. Miljkovic4, B. Jordovic5, Lj. Zivkovic1 1. Faculty of Electronic Engineering, University of Nis, Nis, Serbia 2. Institute of Technical Sciences of SASA, Belgrade, Serbia 3. Faculty of Agriculture, University of Belgrade, Belgrade, Serbia 4. Center for EM, University of Nis, Nis, Serbia 5. Faculty of Technical Sciences, University of Kragujevac, Cacak, Sebia [email protected] Keywords: microstructure, BaTiO3, Rare-earth dopants
In the process of BaTiO3-ceramics consolidation, technological parameters like pressing pressure, initial sample's density, sintering temperature and time, as well as different dopants essentially determine final electrical properties of the ceramics. A slight change of particular consolidation parameter, or the change of dopant's concentration can significantly change the microstructure, thus influencing electrical properties of the speciments. Taking this into account in this article the influence of rare-earth dopants on electrical properties of BaTiO3-ceramics has been investigated. Two types of dopants can be introduced into BaTiO3: large ions of valence 3+ and higher, can be incorporated into Ba2+ positions, while the small ions of valence 5+ and higher, can be incorporated into the Ti4+ sublattice [8-10]. Basically, the extent of the solid solution of a dopant ion in a host structure depends on the site where the dopant ion is incorporated into the host structure, the compensation mechanism and the solid solubility limit [11]. For the rare-earth-ion incorporation into the BaTiO3 lattice, the BaTiO3 defect chemistry mainly depends on the lattice site where the ion is incorporated [12]. It has been shown that the three-valent ions incorporated at the Ba2+ -sites act as donors, which extra donor charge is compensated by ionized Ti vacancies, the threevalent ions incorporated at the Ti4+ -sites act as acceptors which extra charge is compensated by ionized oxygen vacancies, while the ions from the middle of the rareearth series show amphoteric behavior and can occupy both cationic lattice sites in the BaTiO3 structure [11]. As a result, the abnormal grain growth and the formation of deep and shallow traps at grain boundaries influenced by the presence of an acceptor-donor dopant can be observed. Taking this into account in this article, the study of domain structure and microstructure development of BaTiO3 doped with rare-earth dopants has been presented. BaTiO3-ceramics doped with 0.01 up to 0.5 wt. % of Er2O3, Yb2O3 and Ho2O3 were prepared by conventional solid state procedure and sintered up to 1350oC for four hours. Microstructural investigations were carried out, using SEM method and quantitative metallography methods. Domain structure, grain size distribution and porosity of the samples were obtained. Linear intercept measurement method was used for estimating the grain size values and pores volume ratios.
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b)
Figure 1. SEM micrograph of doped BaTiO3 a) 0.5Er-BT and b) 0.5Yb-BT. 100
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Figure 2. Cumulative grain size distribution curves for doped BaTiO3 with a) 0.5 wt% Er2O3, and b) 0.5 wt% Yb2O3 The obtained results enabled establishing, the directions of possible BaTiO3ceramics materials properties prognosis according to the synthesis-structure and structure-property correlations. 1. 2. 3. 4. 5. 6.
1. V.V. Mitić, I. Mitrović, Journal of the European Ceramic Society, Vol. 21 (15), pp. 26932696, 2001. 2. H.M.Chan, M.P.Hamer, D.M.Smyth, J.Am. Ceram. Soc., 69(6) (1986) 507-10. 3. P.W.Rehrig, S.Park, S.Trolier-McKinstry, G.L.Messing, B.Jones, T.Shrout J. Appl. Phys. Vol 86 3, (1999) 1657-1661. 4. D.Makovec, Z.Samardzija M.Drofenik J.Am.Ceram.Soc. 87 [7] 1324-1329 (2004). 5.D. Lu, X. Sun, M. Toda Japanese Journal of Applied Physics Vol. 45, No. 11, 2006, pp. 8782-8788. 6. This research is a part of the project ”Investigation of the relation in triad: synthesisstructure-properties for functional materials” (No.142011G). The authors gratefully acknowledge the financial support of Serbian Ministry for Science for this work.
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SEM and EDS Analysis of BaTiO3 Doped With Yb2O3 and Ho2O3 V. Mitic1,2, V.B. Pavlovic3, V. Paunovic1, M. Miljkovic4, B. Jordovic5, Lj. Zivkovic1 1. Faculty of Electronic Engineering, University of Nis, Nis, Serbia 2. Institute of Technical Sciences of SASA, Belgrade, Serbia 3. Faculty of Agriculture, University of Belgrade, Belgrade, Serbia 4. Center for EM, University of Nis, Nis, Serbia 5. Faculty of Technical Sciences, University of Kragujevac, Cacak, Sebia [email protected] Keywords: SEM, EDS, BaTiO3, Yb2O3, Ho2O3
Because of their high dielectric constant, thermal stability and low losses barium titanate based materials are one of the most common ferroelectrics, with extensive use as a dielectric materials for multilayer ceramic capacitors (MLCCs), embedded capacitance in printed circuit boards, thermal imaging and actuators, dynamic random access memories (DRAM) in integrated circuits [1-4]. The consolidation of ceramics powders on the base of barium-titanate has a great importance, especially from the point of view of further prognosis and properties design of these ceramics. In this article the influence of Yb2O3 and Ho2O3 on microstructure characteristics of BaTiO3-ceramics has been investigated. BaTiO3-ceramics doped with 0.01 up to 0.5 wt. % of Yb2O3 and Ho2O3 were prepared by conventional solid state procedure and sintered up to 1350oC for four hours. Microstructure characterizations for various samples have been carried out by scanning electron microscope of the JEOL-JSM-T20 type, which enables the observation of samples surface by enlarging to 35000 times, with the resolution of 4,5 nm. The application of EDS analysis has been done by energy dispersive spectrometer. EDS-system QX2000S (Oxford Instruments, UK) connected with scanning electron microscope and multichannel analyzer (MCA) is used. Our study has shown that doping with Yb2O3 and Ho2O3 is an effective way to achieve the microstructure control of barium-titanate ceramics. A change of dopant's and their concentration significantly changed the microstructure of the samples. For the lowest concentration, the size of the grains was large (up to 60 μm), but by increasing the dopant concentration the grain size decreased. As a result, for 0,5 % of dopant the average grain size was from 7 to 10 μm. Spiral concentric grain growth which has been noticed only for the samples sintered with 0.5 % of Ho2O3. The formation of the "glassy phase" for the samples sintered with 0.5 % of Yb2O3 indicated that the sintering was done in liquid phase (Figs. 1 a-e). EDS analysis has been shown that for the small concentration of dopants the uniform distribution has been noticed, while the increase of dopant concentration led to the coprecipitation between grains.
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S. Wang, G.O. Dayton J. Am. Ceram. Soc. 82 (10), (1999), 2677–2682. C.Pithan, D.Hennings, R. Waser International Journal of Applied Ceramic Technology 2 (1), (2005), 1–14. B.D.Stojanovic, C.R.Foschini, V.Z.Pejovic, V.B.Pavlovic, J.A.Varela, Journal of the European Ceramic Society 24, (2004), 1467-1471. J. Nowotny, M. Rekas Ceram. Int., 17 (4), (1991), 227-41. This research is a part of the project ”Investigation of the relation in triad: synthesisstructure-properties for functional materials” (No.142011G). The authors gratefully acknowledge the financial support of Serbian Ministry for Science for this work.
a)
c)
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Figure 1. SEM micrograph of doped BaTiO3 sintered at 1320 oC a) 1%Yb-BT, b) 1%Ho-BT, c) 1%Yb-BT, d) 1%Ho-BT These results enabled microstructure control especially intergranular contacts which is important for understanding intergranular capacities and MLCC manufacturing.
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Structure and superconductivity of Pr-Ba-Cu-O crystals prepared by ambient pressure synthesis using citrate pyrolysis method K. Nishio1, T. Isshiki1, T. Shima1 and M. Hagiwara1 1. Kyoto Institute of Technology, Matsugasaki, Sakyo-ku, Kyoto 606-8585, Japan [email protected] Keywords: Cuprate superconductor, Cu-O chain, stacking sequence
Y-base high-Tc cuprate superconductors include three different crystalline phases which are distinguished by each crystalline stacking sequence of layers of Cu-O chains. A single chain layer is placed in a unit cell of YBa2Cu3O7−x (123), while two double chain layers are arranged in that of YBa2Cu4O8 (124). In the Y2Ba4Cu7O15−x (247) phase, alternate single and double chain layers are repeated. It has been recognized that the superconductivity on the CuO2 planes is fully restrained by Pr substitution for Y sites, differently from substitution by other rare earth elements [1]. However, it was found that Pr2Ba4Cu7O15−x (Pr247) showed superconductivity on the Cu-O double chains when the sintered specimen was reduction treated [2], and then Pr-Ba-Cu-O system, especially Pr247, has been attractive with a view to investigating new high-Tc superconductivity. In this work, the stacking patterns found in Pr-Ba-Cu-O sintered material, which was synthesized at ambient pressure condition using citrate pyrolysis method [3], are studied by high-resolution transmission electron microscopy (HRTEM). The Pr-Ba-Cu-O precursor was produced by the citrate pyrolysis method with the nitric acid aqueous solution of source elements [Pr]:[Ba]:[Cu] of the ratio 2:4:7. Pr-BaCu-O powder was obtained by calcining the precursor at 884 °C in atmospheric flowing O2 gas. Such calcined products were shaped into a disk pellet and were sintered at 883 °C in O2. Reduced sample was obtained by heat treatment of the sintered sample at 400 °C in vacuum. Specimens for HRTEM observation were prepared by grinding the sample in ethanol and were dispersed on Cu grids with a carbon-coated holey film for TEM. HRTEM observation was carried out with a JEM-2010/SP (JEOL Ltd.) microscope operated at 200 kV. It was revealed by X-ray diffraction analysis and HRTEM observation that Pr247 was main phase of the obtained material, while a small amount of Pr123 and Pr124 phases was found in addition to BaCuO2 which is one of typical impurities at Pr-Ba-CuO system. Since thermal process for the aimed phases is sensitive to the calcining and sintering conditions [3], it was also found that the material included disorder of the stacking sequence of the single and double Cu-O chains besides Pr123, Pr124 and BaCuO2. Thus it is considered that Pr-Ba-Cu-O material forms heterogeneous particles which are modulated or long-period structure. The particles with the composition between Pr247 phase and Pr123 phase were often observed, as shown in Figure 1. Reduced products including the heterogeneous particles show rather higher Tc of superconductivity than high-purity Pr247 [4]. Besides, it is also confirmed that reduced
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Pr124, which contains neither the single Cu-O chains nor oxygen deficiency, does not show superconductivity. The facts suggest that superconductive property on the double Cu-O chains is affected by the oxygen reduction at the single Cu-O chains in the modulated structure as well as pure Pr247 region. As shown in Figure 2, some of the reduced particles indicate a peculiar contrast on the Cu-O chains, which is interesting to elucidate superconductive property of Pr-Ba-Cu-O system. 1. 2. 3. 4.
S. Hori et al., Physica C 302 (1998), p. 10. M. Matsukawa et al., Physica C 411 (2004), p. 101. M. Hagiwara et al., Physica C 445-448 (2006), p. 111. M. Hagiwara et al., Physica C 463-465 (2007), p. 161.
Figure 1. HRTEM image and corresponding SAED pattern of the sintered sample with long-period structure observed along [100] zone axis. Single and double Cu-O chains are indicated as S and D, respectively.
Figure 2. HRTEM image and corresponding SAED pattern of the reduced sample with peculiar contrasts (indicated by the arrowheads) observed along [010] zone axis.
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Electron Diffuse Scattering in epitaxially grown SrTiO3 thin film F. Pailloux1 and J. Pacaud1 1. Lab PhyMat, UMR 6630, Ave P et M Curie, SP2MI/BP30179, F686962 Chasseneuil Cedex, France [email protected] Keywords: diffuse scattering, diffraction, thin film
SrTiO3 (STO) has generated great interest in recent years due to its large potential for technical applications in electronic devices, due to the similarity of structure between STO and other perovskite materials showing high temperature superconductivity, colossal magnetoresistance or ferro-electromagnetism, it can be used for the design of complex new microelectronic devices involving such materials especially in the field of the spintronic [1]. Moreover, STO is a model system for the perovskite family and the detailed study of its behavior can give insight to the properties of the more complex materials of similar structure. The main requirement for many devices is the growth of a high purity and high structural quality thin film. The perovskite structure is extremely sensitive to the deposition condition and particularly to the temperature and the partial pressure of oxygen. Changes in -deposition conditions may lead to a large deviation of the dielectric properties of thin films from those of bulk materials. The chemistry of defects is often proposed as an explanation of this deviation. Beside the oxygen deficiency, the cation stoichiometry seems to play a major role on the structure and properties of the grown film as it might induce the nucleation of structural defects (dislocation loops, stacking fault, Ruddlesden-Popper faults...). Another important parameter for thin perovskite films is the geometrical constraint imposed by the substrate. Most of the time, perovskite exhibit excellent epitaxy on each other and the films are tied to the substrate so the in-plane parameters are not free to reach their bulk equilibrium values. For materials as sensitive to phase transition as perovskite this effect and the associated relaxation processes can be extremely important for the fine tuning of the physical properties of the film. For perfect crystal, diffuse scattering is mostly inelastic due to phonons, plasmons and other processes. Thermal diffuse scattering or phonon scattering can lead to large effects in the electron diffraction pattern especially near second order phase transition where phonon softening occurs. Additionally, elastic diffuse scattering comes from structural deviations from a periodic lattice [1]. These structural deviations can be defects, partial ordering of otherwise disordered structure or structural fluctuations. By studying diffuse scattering we can obtain information about the crystal imperfections and dynamics which can not be obtained from other characterization methods. The goal of this study is to characterize, through diffuse scattering in electron diffraction, the structure of defects in the epitaxial layers of perovskite structure and
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more specifically its influence on the dynamic of the lattice (lattice vibrations, structural fluctuations and continuous phase transition due to phonon softening) [3]. These properties have to be linked with the anomaly of the ferroelectric properties of such thin films. Different strain states can be probed by changing the substrate material or introducing different buffer layers. 1. 2. 3. 4.
Herranz G., Basletic M., Bibes M., Ranchal R., Hamzic A., Tafra E., Bouzehouane K., Jacquet E., Contour J.P., Barthélémy A., Fert A., Phys. Rev. B, 73, (2006) p. 064403 Krivoglaz M.A., "Diffuse scattering of X-Rays and Neutrons by Fluctuations", Springer Berlin (1994) Wang R., Zhu Y., Shapiro S.M., Phys. Rev. B, 61,(2000) p. 8814 The authors acknowledge Dr K. Bouzehouane from UMR CNRS/Thalès, Palaiseau, France, for providing the samples.
Figure 1. Diffraction patterns along the [061] zone axis for STO film (left) and STO substrate (right). The transmitted beam is out of the picture below the right corner. Black arrows show the sheets of diffuse scattering in the <100>* planes. Red arrow shows the row of low temperature phase spots and the diffuse streak in the [100]* direction.
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Investigation of the hydration of calciumsulfate hemihydrates with different microscopic methods C. Pritzel1, R. Trettin1 1. University Siegen, Institute for Building- and Materials-chemistry, Paul-Bonatz-Str. 9-11, 57068 Siegen, Germany [email protected] Keywords: calciumsulfate, hemihydrate, in situ optical microscopy
Some well-known theories about the progresses when calciumsulfate-hemihydrates react with water exist in literature. The growth of gypsum crystals out of a supersaturated liquid was described by Le Chatelier in 1919 |1|. Another kind of reaction was published by Perederji in 1956 |2| and Eipeltauer in 1960 |3|. They described the crystallisation of gypsum as an inner reaction in the calciumsulfate-subhydrates by an inner transformation of the crystal structure. In literature however there has not been an experimental verification for the second theory so far. Therefore we have investigated the reaction of calciumsulfate-subhydrates with water with optical microscopy. For this we invented a special measuring cell. The results obtained by optical microscopy are also supported by other methods, i.e, heat flow calorimetry and scanning electron microscopy and correlate obviously with each other. Thus we were able to show that the hypothetical mechanism of the reaction which was described by Le Chatelier and that of the reaction published by Periderji and Eipeltauer take place parallelly. The results can be observed in short films and single pictures for α- and β-calciumsulfate-subhydrates. We also tried to influence the reaction by different ways, so we tested a higher reaction temperature or additives |4,5|. To get a higher reaction temperature we used a special hot-stage for microscopy (Metter Toledo Type FP 82 D). In fact we established a special measure-arrangement and used it for the in situ optical microscopy. 1. 2. 3. 4. 5.
M.H. Le Chatelier : Crystalloids against colloids in the theory of cements, Trans. Faraday Soc. 14 (1919) 8. I.A. Perederij: Theorie der Bildung, Erhärtung und Festigkeit von normalem Gips und hochfestem Gips GP, Chem. Techn. 8 (1956) 659. E. Eipeltauer: Erzeugung von krichfesten Hartgipsen, Zem.-Kalk-Gips 6 (1960) 259. Pritzel, C., Trettin, R.: In-situ Untersuchungen zum Reaktionsverlauf von Calciumsulfatsubhydraten mit Wasser, Monographie zur Tagung Bauchemie Erlangen (2004) 252 Pritzel, C., Trettin, R.: Vergleichende Messungen zur Hydratation von Calciumsulfatsubhydraten mit unterschiedlichen Messmethoden, Monographie GDCh – Tagung Bauchemie Karlsruhe 2006 (155-164)
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 561–562, DOI: 10.1007/978-3-540-85226-1_281, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1 (left) + 2 (right) Calciumsulfate in water at the beginning of the reaction and after 180 minutes, inner hydration and growing of gypsum crystals out of a supersaturated liquid
Figure 3 (left) + 4 (right) Calciumsulfate in water at the beginning of the reaction and after 180 minutes, reaction-temperature 40 °C
Figure 5 (left) + 6 (right) Calciumsulfate in water at the beginning of the reaction and after 180 minutes, reaction-temperature 70 °C
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Investigation of holes in calciumsulfate-hemihydrate crystals by different microscopic methods C. Pritzel1, R. Trettin1 1.
University Siegen, Institute for Building- and Materials-chemistry, Paul-Bonatz-Str. 9-11, 57068 Siegen, Germany
[email protected] Keywords: calciumsulfate-hemihydrate
As calciumsulfate-hemihydrate is obtained by recrystallization in autoclaves, some small holes could be found. These holes influence the further crystallization of the hemihydrates. They can only be investigated when the crystals are large enough. We detected those holes with different methods to find out the reason for their construction. We used optical microscopy, scanning electron microscopy (SEM), confocal laser scanning microscopy (CLSM) and atomic force microscopy (AFM). Optical microscopy has been used to detect the holes but it was not possible to get a survey of the complete surface. We used an Olympus BX 61 microscope with the AnalySIS five software. We detected and measured the holes very easily with this software and took some pictures by extended focal imaging (EFI). With this method some pictures are taken in which different parts of the area are in focus. Combining these pictures you obtain a sharp image of the sample. With CLSM it is possible to detect the depth of the holes, but there is less information about the surface interference of the environment. With SEM we also get explicit information about the crystal’s surface around the holes, but we do not get any information about their depth. With the AFM-method we could examine the surrounding structure of the holes closely. We first verified the existence of the holes by optical microscopy. From this investigation we got more information about the material’s structure around the holes. In a next step these structures have been examined by SEM and found out that the structure consists of levels with different hights. To determine the depth of the holes CLSM has been used because the other possible methods cannot be used to inspect thin tunnels. By AFM we found out that the hemihydrates crystals are formed by different layers starting from the centre of the crystal in direction to the holes. From this investigations could be concluded that the holes result from a faster growth of the outer layers. Because of this calcium- or sulfate- ions could not get into the inner layer. If the crystals dry after formation they will break at these weak points.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 563–564, DOI: 10.1007/978-3-540-85226-1_282, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. Hole in calciumsulfatehemihydrate, optical microscopy with EFI-Methode
Figure 2. Hole in calciumsulfatehemihydrate, scanning electron microscopy
Figure 3. Hole in calciumsulfatehemihydrate, confocal laser scanning microscopy
Figure 4. Surface around a hole in calciumsulfate-hemihydrate, atomic force microscopy
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Analytical and high-resolution TEM investigation of Boron-doped CeO2 B. Rahmati1, G. Gregori2, W. Sigle1, C.T. Koch1, P.A. van Aken1, J. Maier2 1. Stuttgart Center for Electron Microscopy, Max Planck Institute for Metals Research, Heisenbergstr. 3, 70569 Stuttgart, Germany 2. Max Planck Institute for Solid State Research, Heisenbergstr. 1, 70569 Stuttgart Germany [email protected] Keywords: B-doped CeO2, analytical TEM, potential mapping
Rare earth-doped ceria counts as a significant ionic conductor that can be used as electrolyte in solid oxide fuel cells [1]. However, the significant electronic conductivity still represents the main limitation for this kind of applications. Electronic conduction affects the space-charge regions adjacent to the grain boundaries. For this reason, there is great interest in understanding what the effect of dopants on the grain-boundary properties is. In the present study, the effect of trivalent boron on the grain-boundaries properties of nano-crystalline ceria was investigated. Boron-doped CeO2 samples were prepared starting from nano-sized CeO2 and B2O3 powders (nominally 10 mol% B2O3) via vibration milling followed by cold isostatic pressing (600 MPa) and sintering in air at 800 °C for 30 minutes. The resulting average grain size determined via XRD was 60 nm. TEM investigations on the grain boundaries were carried out by high-resolution transmission electron microscopy (HRTEM) in a 1250 kV high-voltage microscope (JEOL JEM ARM1250), electron energy-loss spectroscopy (EELS) in a 100 kV STEM (VG HB501UX), and inline holography in a 200 kV FEG-TEM (Zeiss SESAM). Our EELS studies on all investigated grain boundaries showed unambiguously the following results: 1) Segregation of B at grain boundaries (Figure 1). HRTEM studies on the grain boundaries, however, did not reveal any secondary phase at the grain boundaries (see Figure 2 as a representative example). 2) An increase of the Ce-M4,5 white-line intensity ratio (M5/M4) at the grain boundaries compared to the adjacent grains (Figure 3). This indicates a reduction of Ce valency [2] at the grain boundaries. Potential profiles across grain boundaries were determined by inline holography using a recently developed algorithm for reconstructing relative phase shifts from images recorded at largely different defocus values [3]. We will present charge density profiles across grain boundaries extracted from these. Impedance spectroscopy measurements, which are currently on-going will complement the electron microscopy analyses in terms of explaining the defect chemistry of boron-doped ceria and more precisely the compensation mechanism for acceptor-type boron doping.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 565–566, DOI: 10.1007/978-3-540-85226-1_283, © Springer-Verlag Berlin Heidelberg 2008
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Y.H. Cho, P.S. Cho, G. Auchterlonie, Acta Materialia 55 (2007) 4807 L.A..J. Garvie, P.R. Buseck, Journal of Physics and Chemistry 60 (1999) 1943 C.T. Koch, Ultramicroscopy 108 (2008) 141 160
B K-edge
Ce N2,3-edge
Counts .103
counts x 10^3
140
120
Bulk1 100
Bulk 2 Grain boundary
80 190
195
200
205
210
215
eV (eV) Energy loss
220
Figure 1. Comparison of EELS spectra from a grain boundary and the neighboring grains shows clearly the segregation of boron at the grain boundary.
d(111) = 0.3 nm
Figure 2. HRTEM image of a grain boundary in the nanocrystalline boron-doped ceria. 80
Ce M5-edge
70
Ce M4-edge
counts x 10^3
Counts . 103
60 50 40 30
Bulk 1 Bulk 2 Grain boundary
20 10 0
885
890
895
900
905
eV Energy loss (eV)
910
915
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Figure 3. Comparison of Ce-M4,5 EELS spectra from a grain boundary and the neighboring grains shows the increasing of the white-line intensity ratio (M5/M4) at the grain boundary compared to the adjacent grains.
567
Accommodation of the compositional variations in the Ca1-xSrxMnO3-δ (0≤x≤1, 0≤δ≤0.5) system S. de Dios, J. Ramírez-Castellanos, A. Varela, M. Parras, J.M. González Calbet Departamento de Química Inorgánica, Facultad de Ciencias Químicas, Universidad Complutense, 28040-Madrid, España. [email protected]. Keywords: perovskite, non-stoichiometry, short range ordered cubic phases.
AMnO3-δ(A= Ca, Sr) is a complex and phenomenologically rich group of nonstoichiometric compounds related to the perovskite structure. The hexagonal perovskite SrMnO3 crystallizes in the 4H- structure type [1]. This structure based on a mixed hexagonal/cubic close-packing of SrO3 layers (ABAC stacking sequence) consists of pairs of face-sharing octahedra, M2O9 units, linked by corner sharing. CaMnO3 shows a 3C-layer [2] structure with cubic stacking sequence, ABC, where all MnO6 octahedra share corners. Manganites based on the perovskite structure show very interesting magnetic and electronic properties as well as promising technological applications. These properties can be tuned over a wide range through the choice of size and charge of the A-site cations which control structural distortions and the Mn formal valence. Structural and magnetic properties of 4H, orthorhombic and cubic –phases of the Sr1-xCaxMnO3 system have been reported when the anionic sublattice is complete [3]. The existence of a large oxygen deficiency has been reported in SrMnO3-δ [1] and CaMnO3-δ [4]. Recently, the influence of the non-stoichiometry, δ, in CaMnO3-δ on charge transport and magnetic ordering was studied and large CMR effects were reported [5]. Considerable efforts have been made to trying to understand the complex structureproperty relationships associated with perovskite-phases containing large A-cations. A similar effort, however, has not been focused on phases with variable manganese oxidation states. In this sense, we have undertaken the study of Ca1-xSrxMnO3-δ (0≤x≤1 and 0≤δ≤0.50) system with accurate control of oxygen stoichiometry in order to establish the influence of the size of A-site cations, as well as oxygen content in the structure of these materials. In a systematic structural study, different single-phases are obtained, i.e. 4H, orthorhombic (O), tetragonal (T) and cubic (C)-phases as a function of the Ca, Sr cationic ratio and the oxygen content. According to the X-ray and electron diffraction data in nearly stoichiometric Ca1Sr x xMnO3 system a 4H-structural type is stabilized for 0.75≤x≤1, while for 0≤x≤0.5 an orthorhombic CaMnO3-type phase is obtained. In the Ca0.4Sr0.6MnO3-δ samples, orthorhombic CaMnO3-type phase is stabilized for δ=0.05-0.09. Oxygen vacancies form an ordered arrangement in non-stoichiometric Ca1Sr x xMnO3-δ system with δ=0.44-0.50. X-ray and electron diffraction results show that orthorhombic Ca2Mn2O2.5-type is stabilized for 0≤x≤1 (Figure 1). S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 567–568, DOI: 10.1007/978-3-540-85226-1_284, © Springer-Verlag Berlin Heidelberg 2008
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Between the two limits of this series, nearly stoichiometric Ca1-xSrxMnO3 and Ca2Sr x xMn2O2.5-type, cubic (Pm-3m) perovskite phases with disordered oxygen vacancies are obtained for x≥0.5 and δ=0.35-0.50. However, electron diffraction and HRTEM studies show a more complex microstructure due to the existence of different short– range orders of the oxygen vacancies. SAED pattern of Ca0.25Sr0.75MnO2.56 along the [010] zone axis (Figure 2a) shows besides the basic cubic reflections, diffuse scattering due to the presence of short-range order. A careful examination of the contrasts variation in the HRTEM image (Figure 2b) shows nanodomains with fourfold and twofold superstructures along the [101] directions. Accordingly to that, t he nonstoichiometry is accommodated by means of nanodomains formation. Reoxidation of these deficient-oxygen cubic-phases leads to a metaestable singlephase and nearly stoichiometric cubic perovskite, as it will be discussed. a
b
-
101c 020o
202o 101c
[101]
3.2Å
202A
10.4Å
[101]
b
a
3.8 Å
100
[010]
3.8 Å
800
7.6 Å
001
5.2 Å
10.4 Å
10.4 Å
700 600 500
5.2 Å
Figure 1. a) SAED pattern and b) HRTEM along [10-1] zone axis corresponding to Ca0.9Sr0.1MnO2.50. (subindex c and o refers to the cubic perovskite and orthorhombic Ca2Mn2O2.5-types, respectively). Four-fold and two-fold superstructures along the [101]c and [101]c directions, respectively can be observed. Anionic vacancies are ordered leading to an orthorhombic cell: √2ac x 2√2ac x ac. Figure 2. a) SAED pattern along [010] zone axis corresponding to Ca0.25Sr0.75MnO2.56. The reflections are indexed with respect to a cubic perovskite. Nanodomains of the oxygen vacancies ordering can be seen in the HRTEM image (b). Four-fold and two-fold superstructures along the [101] directions are presents.
400 300
1. 2. 3. 4. 5.
0
1
2
nm
3
4
5
T. Negas and R. S. Roth, J. Solid State Chem. 1 (1970), 409. J.B. MacChesney, H.J. William, J.F. Potter and R.C. Sherwood, Phys. Rev. B 164 (1967), 779. O. Chmaissen, B. Dabrowski, S. Kolenisk, J. Mais, D.E. Brown, R. Kruk, P. Prior, B. Pyles and J.D. Jogersen, Phys. Rev. B 64 (2001), 134412. K.R. Poeppelmeier, M.E. Leonowicz, J.C. Scanlon, J.M. Longo and W.B. Yelon, J. Solid State Chem. 45 (1982), 71. Z. Zeng, M. Greenblatt and M. Croft, Phys. Rev. B 59 (1999), 8784.
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Evidence of SrO(SrTiO3)n Ruddlesden-Popper Phases by High Resolution Electron Microscopy and Holography M. Reibold1, E. Gutmann2, A.A. Levin2, A. Rother1, D.C. Meyer2 and H. Lichte1 1. TU Dresden, Institut für Strukturphysik, Triebenberg-Labor, D-01328 Dresden, Zum Triebenberg 50, Germany 2. TU Dresden, Institut für Strukturphysik, D-01062 Dresden, Germany [email protected] Keywords: Ruddlesden Popper Phases, High Resolution Electron Microscopy, Holography
SrTiO3 (STO) is a preferred substrate material for ferroelectric, ferromagnetic or superconducting films. Its special electrical properties strongly depend on deviations from the stoichiometric chemical composition. Due to the reluctance of introducing point defects into the perovskite lattice, a compositional excess of Sr is compensated through the formation of Ruddlesden-Popper Phases (RP) Srn+1TinO3n+1 characterized by insertion of additional SrO planes into the perovskite lattice. These phases belong to a homologous series of so-called crystallographic shear phases [1-3] (Fig.1). It is well known that in STO crystals, exposed to reducing or oxidizing conditions at elevated temperatures, the near-surface structure can change substantially [4-6]. Notably, beside other phenomena cohering with the high mobility of ionic species, the formation of RP phases is observed. They are introduced by accommodation of long range order lattice distortions formed in consequence of local deviations from STO stoichiometry. Furthermore, Meyer et al. [7,8] discovered a reversible tunability of surface-structural distortions in a single-crystalline STO plate under the influence of an external electric field. Besides an ordering of oxygen-vacancies, another possible explanation includes electro-migration of SrO ion complexes and thus reversible formation of RP phases in near-surface regions of the STO crystal. Therefore RP phases are of great fundamental and technological interest. Thin films of Srn+1TinO3n+1 Ruddlesden-Popper phases (n=1,2,3) were grown epitaxially on (001)-oriented single crystalline SrTiO3 substrates [9]. Preparation of the films was performed by chemical solution deposition (CSD) from metalorganic Sr-Ti solutions (rich in Sr) and subsequent conversion into the crystalline state by different thermal treatments in air atmosphere at a maximum temperature of 700°C. Solutions were prepared by a modified Pechini method. For transmission electron microscopy (TEM), thin cross-sectional foils were produced by mechanical grinding, mechanical dimpling and etching by ion milling. The investigations have been carried out with a Philips Tecnai F20 Cs-corr TEM. The existence of different members of the RP phases was approved. Fig. 2 shows an HRTEM image of the sample with a spacing modulation corresponding to the Sr2TiO4 (n = 1) phase region. At the top of the image, the EMS-simulation [10] of the RP-phase is inserted. Fig. 3 shows amplitude and phase image of another RP phase. Simulations of different RP phases were performed to assure the microscopic results.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 569–570, DOI: 10.1007/978-3-540-85226-1_285, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. Structure models of RP phases. Sr atoms are indicated by black spheres. In the center of each oxygen octahedron one finds a Ti atom.
a
Figure 2. HRTEM image. Spacing corresponds to Sr2TiO4 (n = 1). At the top of the image, the EMS-simulation of the RP phases is inserted.
b
Figure 3. Phase (a)- and Amplitude (b) -image of RP-phases (Sr3Ti2O7) 1. 2. 3. 4. 5. 6. 7.
S.N. Ruddlesden, P. Popper, Acta Cryst. 10 (1957) p. 538 S.N. Ruddlesden, P. Popper, Acta Cryst. 11 (1958) p. 54 M.M. Elcombe et al, Acta Cryst. B 47 (1991) p. 305 K.Szot et al, Appl. Phys. A 62 (1996) p. 335 K. Szot, W. Speier, Phys. Rev. B 60 (1999) p. 5909 Y. Liang, D.A. Bonell, J. Am. Ceram. Soc. 78 (1995) p. 2366 D.C. Meyer, A.A. Levin, S. Bayer, A. Gorbunov, W. Pompe, P. Paufler, Appl. Phys. A 80 (2005) p. 515 8. D.C. Meyer, A.A. Levin, T. Leisegang, E. Gutmann, P. Paufler, M. Reibold, W. Pompe, Appl. Phys. A 84 (2006) p. 31 9. E. Gutmann, A.A. Levin, M. Reibold, J. Müller, P. Paufler, D.C. Meyer, Journal of Solid State Chemistry, 177 (2004) p. 1559 10. P.A. Stadelmann, Ultramicroscopy 21 (1987) p. 13
571
New Barium Antimony Aluminates evidenced by TEM techniques R. Retoux, A. Letrouit, M. Hervieu and S. Boudin 1. Laboratoire CRISMAT, UMR CNRS 6508, ENSICAEN, 6 Bd du Maréchal Juin, 14050 Caen cedex France [email protected] Keywords: transmission electron microscopy, aluminate, X ray diffraction, ab initio structure determination
Due to their structural diversity and chemical stability, the aluminates are widely used for various applications as for example sorption, exchange, catalysis, luminescence or transparent conducting properties [1]. The characterization of original and complex frameworks of new aluminates could be limited because most of these compounds are only obtained in the form of polycrystalline samples. After the discovery of the first fullerenoïd oxide [1-3] (complex alkaline earth bismuth aluminates), we decide to investigate the Ba/Sb/Al/O system towards new aluminate frameworks. Characterizations were carried out by combining electron microscopy and X-Ray powder diffraction techniques. The ab initio structural studies of new the phases were realized by determining first the cationic ratios by EDS analyses, the cell parameters and the space group by ED and secondly the crystal structure by X Ray powder diffraction. Then the crystal structure was comforted by comparison between the simulated and experimental HREM images. We present here the new aluminates Ba1.5SbAl5O11.5 [4] (Figures 1,2) and Ba4Sb8Al16O44 (Figure 3). The first Ba1.5+xSbAl5O11.5 oxide crystallizes in an hexagonal cell (P63/mmc, a = b = 5.664 Å, c = 32.159 Å). Its structure consists of a stacking of different layers, noted M, K and P. The M layer is a mixed layer containing SbO6 octahedra and AlO4 tetrahedra. The K layer is a Kagome layer made of AlO6 octahedra. In the P layer, “pillars” are built from Al2O7 groups. The Ba2+ ions lie between the pillars of the P layer and between two successive M layers. The second oxide, Ba4Sb8Al16O44 adopts a Hollandite structure type with SbOx and AlOx polyhedra, forming double rutile chains with oxygen vacancies. The ordering of Sb and Al atoms within the rutile chains induced a tripling of the c parameter. The Ba2+ ions are located inside the tunnels, parallel to the c direction, on two split sites. These characterizations of the present Barium Antimony Aluminates evidenced 2D and 1D structures where ionic mobility, insertions, exchanges can be expected. The studies of such properties will be undertaken. 1. 2. 3. 4.
M. Hervieu, B. Mellenne, R. Retoux, S. Boudin and B. Raveau, Nature Mat., 3 (2004) p.269. an refences therein 1 to 8 R. Retoux, B. Mellenne, S. Boudin, M. Hervieu and B. Raveau, Sol. St. Sci. 7 (2005) p.736. O.I. Lebedev et al, Int. J. of Material research 97 (7) 2006), p. 978-984. A. Letrouit, S. Boudin, R. Retoux and M. Hervieu, Solid State Science in press (2008).
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 571–572, DOI: 10.1007/978-3-540-85226-1_286, © Springer-Verlag Berlin Heidelberg 2008
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c*
[110]*
[ 1 20] *
c*
004
004 112
102
110
100
a*
[110]*
(a)
Hexagonal cell (b) P63/mmc, a=b≈5.6 Å, c≈32.2 Å Figure 1. Ba1.5SbAl5O11.5 : (a) Electron Diffraction patterns (b) experimental, calculated and difference XRPD patterns AlO6 tetrahedron
‘M’ layer
2 nm
‘K’ layer AlO6 octahedron
‘P’ layer
AlO4 tetrahedron
SbO6 octahedron
c
c
Ba2+
(b)
c
b
Ba2+ Ba2+
(a)
P
P
(c) M M Figure 2. Ba1.5SbAl5O11.5 : (a) Projection perpendicularly to the bc plane, and corresponding (b) experimental and calculated (insert) HREM image, (c) Superposition of calculated HREM image and drawing of the structure. b Ba2
a
+
b
c
a
2 nm (Sb/Al)O
(b) (a) Figure 3. Ba4Sb8Al16O44 : (a) Projection of along c, (b) [001] HREM image, superposition of the structure in insert.
573
(Multi-)ferroic domain walls– a combined ab-initio and microscopical investigation A. Rother1, S. Gemming2, D. Geiger1 and N. Spaldin3 1. Institute of Structure Physics, Technische Universitaet Dresden, Germany 2. FZ Dresden-Rossendorf, D-01314 Dresden 3. Materials Research Lab, University of California, Santa Barbara [email protected] Keywords: Multiferroics, DFT, Holography
(Multi-)ferroic materials attracted growing interest during the last decade due to their interesting (multiple)-ordering phenomena and the resulting applications (i.e. nonvolatile memories). Physical properties of boundaries are of particular importance as electronic device dimensions shrink and multiferroic bulk materials have not revealed a sufficient magneto-electric coupling so far. We will combine Density Functional calculations and microscopic techniques to examine basic properties of model boundaries, i.e. BiFeO3 71°/109°/180° domain walls (Fig. 1) and BaTiO3 90°/180° domain walls. The BaTiO3 180° domain wall is considered in the lower energetic parallel and in the higher energetic head-to-head configuration. The DFT calculations are performed within LDA+U on a plane wave basis set. PAW pseudopotentials have been incorporated to represent core states. Both, unit cell dimensions and ion positions have been relaxed to yield minimal energy structures. Transmission Electron Microscopy and in particular Electron Holography are applied to probe electric potential distributions and structure properties at the domain boundaries. Structural changes at the boundary occur due to lattice misfits and reconstruction of electronic orbitals. The thereby produced polarization change at the boundary leads to depolarization fields (Fig. 2). We particularly calculate and investigate such fields and discuss the prospects and problems of depolarization field measurements with TEM techniques. Moreover, the reconfiguration of the band structure at the boundary can lead to completely new physical properties like reduction of the band gap, magnetization change, etc [1, 2]. We calculate boundary band structures predicting ferromagnetic BiFeO3 71°/180° domain walls, band gap reduction, etc. We will discuss the charge modulation at the boundaries and calculate possible surface charges, in particular at the BaTiO3 180° head-to-head boundary. 1. 2.
J. Rondinelli, M. Stengel and N. Spaldin, arXiv:0706.2199 T. Zhao et al., Nature Materials, Vol 5, (2006)
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 573–574, DOI: 10.1007/978-3-540-85226-1_287, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. Fully relaxed structure of 109° BiFeO3 domain boundary with Bi-O termination calculated with DFT
electrostatic potential in eV
0.4 0.2 0 -0.2
-0.4
Figure 2. Potential jump / depolarization field across the 109° BiFeO3 domain boundary depicted in Fig. 1
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Diagnostic of Li battery cathode by EELS, first principles calculation and spectrum-imaging with multi-variate analysis K. Tatsumi1, Y. Sasano1, S. Muto1, T. Sasaki2, Y. Takeuchi2, K. Horibuchi2, Y. Ukyo2 1. Department of Materials, Physics and Energy Engineering, Nagoya University, Nagoya, 456-8587, Japan 2. TOYOTA Central R&D Labs., Nagakute, Aichi, 480-1192, Japan [email protected] Keywords: LiNiO2, degradation, (S)TEM-EELS
Although LiNiO2 based materials are expected for cathodes of Li rechargeable batteries in high-power use, the electrochemical properties of the cathodes are degraded by charge-discharge cycles at elevated temperatures[1]. Small amounts of Al and Mg are doped so as to suppress the degradation. The purpose of this study is a cathode diagnostics by two different approaches. One is the local environment analysis of the dilute dopants by TEM-EELS and first principles calculations. The other is STEMEELS imaging with a multi-variate analysis for spatial distribution of the degraded areas in the cathode active materials. The samples were LiNi0.8Co0.15Al0.05(Mg0.05)O2 cathode materials, taking a form of secondary particles composed of aggregating primary grains of approximately 1 μm size (Figure.1-a). After the charge-discharge cycling tests on battery cells, the cells were disassembled to take out the cathode. K-shell ELNES of the dopants Al and Mg were measured from the ion-milled sample. Their theoretical spectra were calculated by the first principles APW+lo band method. In order to image electrochemically degraded areas included over an entire secondary particle, samples prepared by FIB were analyzed by STEM-EELS. Figure. 1-b shows the typical Ni-L2,3 and O-K ELNES of the sample. Ni L2/L3 relative intensity ratio is known to become higher as Ni is more oxidized. The broken line spectrum is responsible for Ni2+, while the solid line corresponds to Ni3+ in LiNiO2[2,3]. The O-K spectra differ correspondingly. Because Ni2+ is electrochemically inactive, the areas responsible for the broken line spectra should be degraded. Regarding the Ni L2/L3 relative intensity ratio as a measure of the degradation, the dopants spectra were collected at the normal and anomalous areas. Representative Mgand Al-K spectra are shown in Figures. 2-a and b, respectively. We use the calculated spectra as theoretical fingerprints (Fig. 2-c and d). Comparing the experimental spectra with the theoretical, in the normal LiNiO2 area, we find that both Mg and Al mainly occupy the Ni site. In the anomalous areas, the chemical state of Mg is estimated as Mg in NiO. Similarly, Al in degraded area 1 corresponds to Al in NiO. The spectrum profile of degraded area 2 is consistent with the theoretical one of Al in LiAlO2, which has the same crystal structure as LiNiO2.
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Our EDX and x-ray absorption spectroscopy analyses on the dopants indicated that the charge-discharge cycles decreased the Mg concentration at least at the surface of the secondary particles. Our first principles calculation for the dopants in LiNiO2 indicated that Al-O bonds are much stronger than Mg-O, while Mg reinforced the Ni-O bonds of the nearest neighbor Ni. Combining these findings to the fact that some Al atoms have chemical states similar to LiAlO2, we can see different characters of Mg and Al, that is, the doped Mg would be eluted out from the cathode active material while the Al would stay at the Ni site in LiNiO2, depending on their local chemical bonding. In the presentation, we will discuss the degradation process based on the dopants chemical bonding and the degraded areas distributed in the secondary particle, which we revealed for the first time by a multi-variate analysis on the STEM-EELS results. 1. 2. 3.
H. Kondo et al., J. Power Sources, 174 (2007) 1131. D.P. Abraham et al., J. Electrochem. Soc., 150 (2003) A1450. Y. Koyama et al., J. Phys. Chem. B, 109 (2005) 10749.
Figure 1. STEM dark field image of the cathode material prepared by FIB (a), Typical O-K and Ni-L2,3 ELNES of the cathode (b).
Figure 2. Experimental Mg-K (a) and Al-K (b) ELNES and theoretical ones for possible chemical states of Mg (c) and Al (d)
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Local electronic structure analysis on Ce3+-containing materials by TEM-EELS and first principles calculations K. Tatsumi1, I. Nishida1 and S. Muto1 1. Department of Materials, Physics and Energy Engineering, Nagoya University, Nagoya, 456-8587, Japan [email protected] Keywords: local atomic and electronic structure, Ce3+, ELNES, first principles calculation
A Ce 4f electron in Ce3+-containing materials exhibits specific properties such as luminescence and fast oxygen-absorption / desorption. In their EEL spectra, the Ce M4/M5 relative peak intensity ratios are known as a measure of oxidation states of Ce, although fine structures in the M4 and M5 edge are difficult to be understood by electronic structure calculations within the one-electron approximation. The spectral structure of anionic edge (e.g. O-K and F-K ELNES) is additional useful information to understand the electronic and atomic structures of the materials because it reflects generally the unoccupied electronic density of states (DOS) that could be estimated by the calculations. This study investigates the electronic/atomic structures of (1) typical materials containing Ce3+ (CeF3 and Ce2O3) and (2) ceria-zirconia ordered phases which are used as a co-catalyst for automobile applications. (1) CeF3 and Ce2O3: Figure 1 compares between the experimental and theoretical anion K-shell spectra. In both compounds, the spectral shape of the GGA+U result is more consistent with that of the experimental. As indicated in the figure, the energy width of the main peaks of the oxide spectrum is larger than that of the fluoride. This energy region reflects the width of Ce d bands. From the viewpoint of atomic orbital interaction, it is found that 1) the energy width is formed by atomic orbital interaction among atoms within an approximately 7 Å area (100 atoms), 2) the difference in the energy width between the compounds is due to the difference in the antibonding interaction between Ce 5d and anion 2p; CeF3 in which Ce is more isotropically coordinated by anions and the interaction is homogeneous over different directions of the 5d orbitals, has the narrower Ce 5d energy band. These findings suggest that the bandwidths of Ce f-d luminescence among different Ce3+ containing materials could be clarified systematically according to the local atomic arrangements around Ce. (2) Ceria-zirconia ordered phases: Figure 2 shows the O-K spectra of Ce2Zr2O7, which is the reduced phase in the ceria-zirconia system (Ce2Zr2OX X=7, 7.5 and 8). In this phase, oxygen vacancies are ordered and accordingly O ions are displaced from the ideal positions based on the fluorite structure. As shown in Fig.2-d, O ions are displaced toward Zr and away from Ce. Between the theoretical spectra (Fig.2-b) of the ideal and displaced models, the spectral profile of the displaced model is more consistent with that of the experimental. The partial DOS (Fig. 2-c) indicated that the electronic structures ranging over about 10 eV from the Fermi level are significantly different between the models. The oxygen displacements split Zr(Ce) d DOS peaks in larger (smaller) energies, which suggests that the displaced model has stronger Zr-O bonds and nearly free Ce ions. S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 577–578, DOI: 10.1007/978-3-540-85226-1_289, © Springer-Verlag Berlin Heidelberg 2008
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1. 2.
L.A.J. Garvie et al., J. Phys. Chem. Solids, 60 (1999) p. 1943. S. Muto et al., J. Electron Microscopy, 55 (2006) p. 225.
Figure 1. Experimental and theoretical F-K (a) and O-K (b) spectra. The experimental F-K spectrum is restored by the procedure reported in ref. [2]. The spectra were aligned so that the main peaks are at the same energies.
Figure 2. Experimental (a) and theoretical (b) O-K spectra of Ce2Zr2O7, corresponding partial density of states of the cations (c) and oxygen displacement in this compound (d).
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Local chemical inhomogeneities in NaNb1-xTaxO3 as observed by HRTEM and HAADF-STEM A. Torres-Pardo1, E. García-González1, J.M. González-Calbet1, F. Krumeich2 and R. Nesper2 1. Departamento de Química Inorgánica I, Facultad de Ciencias Químicas, Universidad Complutense de Madrid, 28040-Madrid. Spain. 2. Laboratory of Inorganic Chemistry, ETH Zurich, CH-8093 Zurich, Switzerland [email protected] Keywords: HRTEM, HAADF-STEM, ferroelectric materials, perovskite-solid solution
Sodium niobate NaNbO3, crystallizing with a perovskite-type structure, is an interesting material both because of its structural phase transitions as well as for its electric properties ranging from antiferroelectric to paraelectric behaviour [1]. Compositional variations in the sublattice of both A and B cations have been widely studied in order to stabilize a ferroelectric phase at room temperature [2,3]. The substitution of Ta by Nb leads to the formation of ferroelectric materials for Ta-rich compositions. Remarkably, a noticeable discontinuity of the electric properties occurs in the range 0.4 ≤ x ≤ 0.6 for NaNb1-xTaxO3 and no explanation has been given to justify such behavior [4]. To obtain a deeper insight into the complicated relationship between structure and properties in this system, we prepared and reinvestigated the solid solution NaNb1-xTaxO3 (0 ≤ x ≤ 1.0). X-ray powder diffraction data show a gradual structural evolution between the two limiting compositions. However, the investigations performed by means of SAED and HRTEM (performed on a JEOL JEM300FEG electron microscope) reveal a complex microstructure constituted by three dimensional structural domains of a √2ac×2ac×√2ac unit cell (ac refers to the basic cubic perovskite cell). In addition, recurrent antiphase boundaries-type (APB) defects are observed in the interval 0.4 ≤ x ≤ 0.6 along the [10 1 ] zone axis showing a trend towards an ordered distribution (Figure 1b). The APBs occurring through a translation along ½ [010] lead to a splitting of 0k0 diffraction maxima parallel to the [100]c direction (Figure 1a). Defects seem to maintain both their relative concentration as well as their periodicity in the whole range of composition. The average chemical composition of the defects does not deviate from the nominal one as observed from energy-dispersive spectrometry (EDS) X-ray nanoanalysis. Microstructure characterization by means of HAADF-STEM (performed on a Tecnai F30 electron microscope) has revealed a trend to an ordered Nb/Ta distribution in the B position of the perovskite structure where Nb-rich planes alternate with Ta-rich planes along the [100]c direction of the perovskite structure (Figure 2). All APBs appear to be associated with this ordered cationic distribution and they are not present in crystals with a random Nb/Ta distribution in the B sublattice. These observations indicate that compositional variations are responsible for the formation of structural defects observed in the NaNb1-xTaxO3 (0.4 ≤ x ≤ 0.6) system.
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I. Lefkowitz, K. Lukaszewicz and H. D. Megaw, Acta Cryst, 20 (1966), p. 670. I. P. Raevski and S. A Prosandeev, J. Phys. Chem. Sol. 63 (2002), p. 1939. R. Jiménez, M. L. Sanjuán and B. Jiménez, J. Phys.: Condens. Matter. 16 (2004) p. 7493. H. Ywasaky, Rev. Electr. Commun. Lab.12 (1964), p. 469. We kindly acknowledge the facilities of the Centro de Microscopía Electrónica Luis Brú (UCM) and the Electron Microscopy Center ETH Zurich (EMEZ).
Figure 1. (a) SAED pattern of a crystal of NaNb0.4Ta0.6O3 along [10 1 ]. (b) Corresponding HRTEM micrograph, with the Fourier Transform shown as inset.
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Figure 2. (a) HRTEM and (b) HAADF-HRSTEM image of a NaNb0.4Ta0.6O3 crystal along [10 1 ]. Corresponding Fourier Transforms and the intensity profile from the dotted marked line are shown as insets.
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The influence of lanthanum doping on the structure of PbZr0.9Ti0.1O3 ceramics R. Villaurrutia1, I. MacLaren1, A. Peláiz-Barranco2 1. Department of Physics and Astronomy, University of Glasgow, Glasgow G12 8QQ, UK 2. Facultad de Física-Instituto de Ciencia y Tecnología de Materiales, Universidad de la Habana, San Lázaro y L, Vedado, La Habana 10400, Cuba [email protected] Keywords: Ferroelectric, ceramics, incommensurate
Ferroelectric materials are widely used in applications ranging from piezoelectric through pyroelectric to electro-optic devices and further miniaturisation of ferroelectric devices, for instance in ferroelectric random access memories requires a better understanding of the nanostructure of these materials. Recently, we have shown that orientation mapping from diffraction allows the study of local crystallography and stress concentrations in tetragonal lead zirconate titanate (PZT) [1,2]. Nevertheless, rhombohedral PZT has been less studied, despite its commercial significance [3]. Nominally rhombohedral Lanthanum modified Lead Zirconate Titanate (PLZT) Pb1xLax(Zr1-yTiy)1-x/4O3 has been investigated using Transmission Electron Microscopy including conventional imaging and selected area and convergent beam diffraction techniques. Previous studies have reported the coexistence of ferroelectric (FE) and antiferroelectric (AFE) phases in Zr-rich PLZT with Zr/Ti ratios and La amounts close to the (AFE-FE) phase boundary; the antiferroelectric domains being nanoscale regions within the conventional ferroelectric domains [4]. The present work concentrates on more detailed studies of this domain structure using recent advances in automated analysis of diffraction patterns. Figures 1a and b are dark-field images showing the coexistence of conventional domains and nanoscale domains for the PLZT 2/90/10 composition. The nanoscale domains are streaked in one direction in each domain. Trace analysis of the streaking directions show that these are consistent with a layered structure on {110}Cubic planes. In both images, insets are provided of selected area diffraction patterns recorded at the same orientation, and satellite reflections are associated with most major reflections lying in <110>Cubic directions. The spacing of these reflections corresponds to a spacing of 26 ± 2 Å and 23 ± 2 Å for the two diffraction patterns shown, corresponding to about 8 or 9 (110) spacings, whether this periodicity is truly incommensurate with the lattice is unclear. Similar observations have been made for PLZT 3/90/10 and 4/90/10. Figure 2 shows a bright field image of a PLZT 5/90/10 material and whilst a conventional domain structure on the scale of hundreds of nm, this shows no sign of the streaked structure found at lower La contents, and no sign of this phase could be found in either images or diffraction patterns. This composition is clearly in a different phase field, which broadly accords with the results of Dai et al. [4].
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I. MacLaren, L.A. Schmitt, H. Fuess, H. Kungl, and M.J. Hoffmann, J. Appl. Phys., 97, (2005) 094102. M.U. Farooq, R. Villaurrutia, I. MacLaren, H. Kungl, M.J. Hoffmann, J.J Fundenberger and E. Bouzy, Journal of Microscopy, (2008) in press. J. Ricote, R. W. Whatmore and D.J. Barber, J. Phys.: Condens. Matter, 12 (2000) 323-337. X.H. Dai, Z. Xu, J.-F. Li, and D. Viehland, J. Mater. Res., 11, 3, (1996) 626-638.
Figure 1. Dark field images from the PLZT 2/90/10 ceramic showing clear streaking in some domains, with insets showing the spot splitting arising from the incommensurate phases.
Figure 2. Bright field image of PLZT 5/90/10 showing a conventional domain structure and no sign of nanoscale domains.
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Anomalous absorption of electrons during electron diffraction on BaTiO3 single crystals near phase transition at 120°C A. Wall FH Wiesbaden, FB Physik, Am Brückweg 26, 65428 Rüsselsheim, Germany [email protected] Keywords: Bariumtitanat, phase transition, electron diffraction
Apart from the domain structures that were discovered on BaTiO3 single crystals by means of transmission electron microscopy (TEM) [1, 2, 3], there are a couple of other effects we know very little about. Some of these effects that require further research are: Transition structures [4], or optical radiation which accompanies the diffraction of electrons close to the temperature of phase transition as well as far below, [5,6]. In this article we shall present results of experiments on the anomalous absorption of electrons during electron diffraction on BaTiO3 single crystals at 50keV. The BaTiO3 single crystals were grown according to the Remeika method [7]. The crystals were then corroded by phosphoric acid at temperatures of about 200°C until they reached their adequate thickness (300-500Ǻ). A typical diffraction pattern of BaTiO3 single crystals far below phase transition is shown on Figure No 1. Due to the electric field of the ferroelectric material [3] the photograph is slightly asymmetrical. The white spots on this photo are caused by the surplus of electrons. In some cases, in addition to the typical maxima, dark spots will also result when the crystal is being warmed up by electrons (Fig.2). Further enlargement of the negative patterns will make it possible to carry out more detailed studies on this phenomenon observed in connection with electron diffraction (Fig.3, 4, 5). Here, the dark spots which indicate the absence of electrons become white whereas the diffraction maxima look dark as a result of a surplus of electrons on the negative patterns. There are different kinds of electron diffraction leading to the diffraction maxima and to a development of dark spots. This is made obvious by photometric results (Fig.6). A surplus of electrons near the maximum of diffraction is to be seen on Illustration 6. The shape of the maximum is typical of electron diffraction when electrons are dispersed on atoms of the solid state body. The changing intensity within the spot caused by anomalous absorption (Fig.6) looks very similar to the arc sine distribution. It can be presumed that when it comes to phase transition of BaTiO3 during the transformation of the crystal lattice, the accelerated electrons as well as the photons penetrate deeper into the crystal lattice than they do when metals or dielectric material are involved. This is the reason for the anomalous absorption of electrons. 1. 2.
M. Tanaka, G. Honjo Journal of the Physical Society of Japan 19, 954, (1964). H. Pfisterer, W. Liesk, Physikalische Blätter 24, 488 (1968)
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H. Lichte, M. Reibold, K. Brand, M. Lehmann, Ultramicroscopy 93 (2002) 199-212 E.W. Bursian, A.B. Wall, N.N. Trunov, Sov.Phys. Solid State, 19(197) 1108 A. Wall, Optik 90, 187, (1992) A. Wall, Microscopy and Microanalysis, Vol. 13, Supplement S03, Sep. 2007, pp 374-375 J. R. Remeika. J. Amer. Chem. Soc., 76, 940 (1954)
Figure 1. Typical diffraction pattern of a Figure 4. A greatly enlarged white spot BaTiO3 mono crystal (positive Pattern) (negative pattern).
Figure 2. Diffraction pattern near phase transition (120°C). The position of the Bragg maxima is irregular. Little dark spots can clearly be identified (positive pattern).
Figure 5. Illustration of the diffraction maximum with a satellite. The white spot indicating the absence of electrons can be seen on the left of the lower part (negative pattern). 250
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Template-assisted synthesis and characterization of SrTiO3 nanostructures K. Žagar, S. Šturm and M. Čeh Department for Nanostructured Materials, Jožef Stefan Institute, Jamova cesta 39, 1000 Ljubljana, Slovenia [email protected] Keywords: SrTiO3, nanostructures, sol-gel electrophoretic deposition, transmission electron microscopy
Complex metal oxides with perovskite structure are important materials for various applications in microelectronics industry and technology. They can be used as sensors, catalysts and composites with defined electrical properties. For many applications the physical properties of these perovskite ceramics strongly depend on their surface/volume ratio. Processing of these materials in the form of nanostructures (nanorods, nanotubes) leads to an increase in the surface area compared to their bulk counterparts [1, 2]. Nanostructures of various complex metal oxides have been synthesized by various methods, among which sol-gel template-assisted method provides a versatile technique for synthesizing one-dimensional nanostructured materials [3]. In our work we report on the synthesis of SrTiO3 nanostructures by sol-gel electrophoretic deposition into template membranes. Two types of template membranes were used: track-etched hydrophilic polycarbonate (PC) membranes and anodic aluminum oxide (AAO) membranes, both with pore diameters of 200 nm and thickness of 10-25 μm. The template membrane was attached to aluminum working electrode while Pt mesh electrode was used as a counter electrode. For electrophoretic deposition of the sol into porous templates the potential of 30 V was applied between both electrodes for 30 min. After the deposition, the template membranes were first dried at 200 °C for 12 h and then annealed at 700 °C for 1 h or at 800 °C also for 1 h. Such heating procedure was conducted in order to make the nanostructures dense and crystalline. During heating, the polycarbonate membrane was burnt off while the anodic aluminum oxide membranes were removed after annealing in 6 M NaOH. Obtained nanostructures were characterized by scanning and transmission electron microscopy (SEM, TEM). By using SEM we found that the nanostructures grown in a PC or in an AAO membrane were rod-like in shape having a uniform diameter throughout their entire length. EDXS analysis confirmed that the composition of the observed nanostructures corresponded to stoichiometric SrTiO3. We also observed that higher potential led to a thin SrTiO3 layer formation on the membrane surface after the pores were filled. Figure 1 shows TEM images of a single SrTiO3 nanostructure grown in a PC (Fig. 1a) and in an AAO (Fig. 1b) membrane. While the SrTiO3 nanostructures formed in the PC template appear to be nanorods, the SrTiO3 nanostructures formed in AAO templates may well be nanotubes since the outer edges exhibit much stronger contrast. Both nanostructures are dense and continuous with a diameter ranging from S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 585–586, DOI: 10.1007/978-3-540-85226-1_293, © Springer-Verlag Berlin Heidelberg 2008
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100 to 250 nm and are polycrystalline in nature. Figure 2a shows an enlarged region of such a polycrystalline nanostructure with the grain size raging from 25 up to 50 nm. The comparison between the experimental electron diffraction pattern and the calculated pattern confirms that the grains correspond to cubic SrTiO3 (Fig. 2b). High-resolution TEM showed perfect perovskite pattern with no observable structural defects (Fig. 2c). It was concluded that the electrophoretic deposition of sols into porous templates is a useful technique for preparation of one-dimensional perovskite nanostructures.
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Figure 1. TEM images of an individual SrTiO3 nanostructures grown in PC (a) and AAO (b) membrane.
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Figure 2. (a) TEM image of polycrystalline SrTiO3 nanostructure (grown in AAO membrane) with (b) the corresponding electron diffraction pattern and calculated pattern. (c) HRTEM image of SrTiO3 grain within the nanostructure. 1. S.J. Limmer et.al., J. Mater. Sci. 39 (2004), p. 895. 2. Y. Lin et.al., Appl. Phys. A 78 (2004), p. 1197. 3. S. Singh, S.B. Krupanidhi, Phys. Lett. A 367 (2007), p. 356. Acknowledgements: This work was financially supported by the Ministry of Higher Education, Science and Technology of the Republic of Slovenia and from the European Union under the Framework 6 program under a contract for an Integrated Infrastructure Initiative. Reference 026019 ESTEEM.
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(S)TEM/EELS characterisation of a multilayer C/Cr PVD coating Z. Zhou1, W.M. Rainforth1, M. Gass2, A. Bleloch2, and P.Eh. Hovsepian3 1. Dept of Engineering Materials, University of Sheffield, Sheffield, S1 3JD, UK 2. SuperSTEM at Daresbury Laboratory, Daresbury, Cheshire, WA4 4AD, UK 3. MERI, Sheffield Hallam University, Sheffield, S1 1WB, 1UK [email protected] Keywords: C/Cr PVD coating, EFTEM, STEM/EELS
Carbon based coatings have attracted considerable interest in a wide range of industries due to their excellent tribological properties. Nanoscale hydrogen free C/Cr coatings have been produced by unbalanced magnetron (UBM) sputtering [1]. Their structure and properties are strongly influenced by the deposition bias voltage [1-2]. The multilayer structure at -350V substrate bias is believed to be a result of phase separation and formation of a self-organised layered nanostructure in conditions of high ion irradiation and elemental diffusion [2]. Composition profile across the alternating layers is the key to understanding the selforganised growth mechanism. We used a dedicated Cs-corrected STEM coupled with electron energy loss spectroscopy (EELS), which is capable of sub angstrom high angle annular dark field (HAADF) resolution. Line scans across the multilayers were performed using a 0.1nm probe. Additional elemental mapping of the nanocomposite structure was obtained using a 2010FEGTEM equipped with Gatan GIF 2000. Figure 1a shows a bright field TEM image of the multilayer and its selected area diffraction pattern (inset). The coating is composed of a 1.00±0.02μm thick multilayer C/Cr on top of a 250±20 nm thick CrN base layer grown on a stainless steel substrate. HREM micrograph in Figure 1b shows that the multilayer consists of nanocrystalline matrix with amorphous inclusions. Figure 2 gives a typical area of the multilayer and its corresponding energy filtered maps by using C K and Cr L2,3 electron energy loss edges. It revealed that the matrix is Cr-rich and the inclusions are C-rich. The C-rich inclusions formed in layers alternated by the Cr-rich layers. Figure3a shows a STEM bright field image at the bottom of the multilayer with a box inset indicating where a spectrum image was recorded. The quantitative EELS line profiles of C, N and Cr processed from the spectrum image is displayed in Figure 3b. It suggests the Cr-rich matrix have some C and vice versa. Carbon is evenly distributed across the Cr carbide layer between two carbon inclusions. EELS near edge fine structure of C K of the C-rich inclusion and Crrich matrix are given in Figure 3c. There is no significant difference between C-rich inclusions and Cr-rich matrix except the ratio of sp3/sp2. C-rich inclusions have higher ratio, as shown in Figure 3d. 1. 2.
P.Eh. Hovsepian, Y.N. Kok, A.P. Ehiasarian, A. Erdemir, J.-G. Wen, I. Petrov. Thin Solid Films 447-448 (2004), p.7. P.Eh. Hovsepian, Y.N. Kok, A.P. Ehiasarian, R. Haasch, J.G. Wen, I. Petrov. Surf. & Coat. Technol., 200(2005) p.1572.
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We kindly acknowledge Engineering Physical Sciences Research Council, UK for funding.
Figure 1. (a) TEM bright field image of the multilayer C/Cr coating and selected area diffraction pattern (inset). (b) HREM micrograph of the multilayer shows amorphous inclusions in a nanocrystalline matrix.
Figure 2. Bright field TEM image of an area of the multilayer and its corresponding EFTEM maps using C K and Cr L2,3 edges showing the inclusions are carbon rich and matrix is Cr rich. 70000
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Figure 3. (a) Bright field STEM image of an area at bottom of multilayer. The box (inset) indicates the location of EEL spectrum image. (b) Line profiles of C, N and Cr intensity. (c) EELS point analysis of carbon rich inclusion and Cr rich matrix. (d) Line profile of sp3/sp2 derived from C K.
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High resolution imaging of magnetic structures in a TEM – what is possible? J. Zweck Electron microscopy lab, Physics Faculty, University of Regensburg, 93040 Regensburg, FRG [email protected] Keywords: Lorentz microscopy, magnetic imaging, differential phase microscopy
Enormous efforts are under way to create devices which use not only the charge of electrons as in conventional electronics but also the long neglected spin properties for a future “spintronics” world. Such magnetic devices are potentially faster, smaller, multifunctional, need less energy to operate and offer non-volatility in the case of power failure which would be most useful for memory devices. As magnetism is a collective phenomenon, it is obvious that once the number of interacting atoms falls below a certain limit (roughly some 106 atoms), the magnetic properties fade away. For possible applications it is therefore absolutely necessary to be able to investigate the magnetic properties of tiny magnetic particles (i.e. down to or smaller than approx. 100 x 100 x 10 nm3) with high resolution. The term 'resolution' has here to be understood in a double meaning, namely as conventional lateral resolution and as magnetic resolution, i.e. what are the smallest detectable changes in magnetic induction. The use of electron microscopy for the investigation of magnetic specimens may seem to be counterproductive at least, if not senseless, as one has to expect that the high magnetic fields of objective lenses simply wipe out all micromagnetic information within the specimen. This problem can be overcome with the use of a so-called Lorentz lens, a long focal-length lens which acts as a remote objective lens. The standard objective lens is turned off, while the Lorentz lens uses a much smaller field in a larger distance from the specimen, leaving the specimen area virtually field free with a remaining lateral resolution of 2 nm. In combination with a dedicated specimen holder, which allows the in-situ creation and precise control of magnetic fields at the specimen's location, the micromagnetic properties can be precisely measured even from from individual particles. In this contribution, an overview will be given both on instrumental requirements and possibilities and on the three dominantly used techniques for micromagnetic investigations in a TEM: the Fresnel imaging mode, differential phase contrast imaging and electron holography on magnetic specimens. Fresnel imaging [1] is a rather simple and easy-to-use technique for quick specimen inspection with inherent advantages and disadvantages. It is an out-of-focus technique which allows to tune the sensitivity for magnetic features by changing the defocus. However, the defocus causes an image blur and – for patterned media – edge fringes which severely obstruct the wanted information. Besides, the contrast formation is highly non-linear and quantitatively difficult to interpret.
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Higher resolution can be achieved by differential phase contrast imaging [1], a STEM technique where a small electron probe is scanned across the specimen. Depending on the local induction, the electron beam becomes deflected and the deflection is monitored by a position sensitive device. The lateral resolution in this mode is limited by the size of the electron probe on the specimen, which (FEI Tecnai F30ST Regensburg special2) is as low as 5 nm. This is sufficient to measure the hysteretic behaviour of small magnetic particles down to a diameter of 100 nm and less. Another high resolution technique is available by electron holography [2], where the phase shift of the electron wave due to interaction with the magnetic flux is measured and evaluated quantitatively. A representative example of a hysteresis loop from electron holograms is given in Figure 1. The ability of electron microscopy to image and in-situ manipulate the magnetic properties of individual magnetic particles is demonstrated by many examples. They focus on the role of exact specimen geometry on magnetic behaviour. It will be shown, how magnetic particles behave differently depending on their size and shape. This can – in turn – be used to specifically design the magnetic properties according to one's needs, but also lead to undesirable behaviour. One is even able to trap magnetic vortices in artificially created holes and to switch the vortices between adjacent holes, leading to multistable states which can act as a memory [3]. 1. 2. 3.
J. N. Chapman, J. Phys. D: Appl. Phys. 17, (1984), p. 623-647 R. E. Dunin-Borkowski, M. R. McCartney, D. J. Smith, and S. S. P. Parkin, Ultramicroscopy 74, (1998), p. 61-73. M. Rahm, J. Stahl, W. Wegscheider, and D. Weiss, Appl. Phys. Lett. 85, 9, (2004), p.15531555.
Figure 1. Middle: Example of a hysteresis loop extracted from electron holography data. Upper and lower row of images: a: shadow image of specimen, a Ni81Fe19 alloy of 15 nm thickness. b - j: reconstructed electron holograms showing the micromagnetic configuration within and outside the specimen for various external applied fields. The letters on the images correspond to those in the hysteresis loop, indicating the configurations for several points of the loop.
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Phase segregation leading to spontaneous outcropping of (Sr,La)Ox dots in La1-xSrxMnO3 films P. Abellan1, F. Sandiumenge1, C. Moreno1, M.J. Casanove2, T. Puig1 and X. Obradors1 1. Institut de Ciència de Materials de Barcelona, CSIC, 08193 Bellaterra, Spain 2. CEMES, 29 rue Jeanne Marvig, BP 94347, 31055 Toulouse Cedex 4, France [email protected] Keywords: LSMO films, MOD, phase segregation, epitaxy
La1-xSrxMnO3 (LSMO) with x=0.3-0.4 exhibits metallic ferromagnetism at room temperature and offers good lattice mismatch with YBCO. Such properties bring LSMO an interesting potential as template for the growth of epitaxial superconducting YBCO. Recently, we have found that (Sr,La)Ox dots can be generated by simple thermal processing while keeping the structural integrity of the films unaltered [1]. This finding offers an interesting scenario for the investigation of the effect of structural disturbances and ferromagnetism on the flux pinning capabilities of YBCO films. Furthermore, these templates are prepared in a one step process following a low cost, easily scalable growth route. LSMO films of high crystallographic perfection and high quality magnetic properties have been grown by MOD on STO(001) substrates. After a suitable thermal treatment, spontaneous phase segregation occurs which results on self-assembled (Sr,La)Ox superficial dots. Formation of these islands is accompanied by the segregation of a polytypoid phase of the manganite structure inside the film, Figure 1. The appearance of these inclusions is proposed to occur for stoichiometric balance of the system. Strikingly, this complex phase-segregation process do not disturb the LSMO structural integrity, which keeps a ferromagnetic transition above room temperature, Tc=360K. The LSMO films are rhombohedral, around 24 nm thick, and present very low surface roughness. The cation distribution has been investigated by energy filtered TEM (EFTEM) and inter-diffusion effects between the LSMO film and STO substrate have been ruled out. Diffraction contrast imaging of planar view foils revealed the presence of the two twin variants along the (010) and (100) planes normal to the interface. At some places the interface also revealed misfit dislocations with Burgers vectors b=a[100], a being the lattice parameter of the pseudo-cubic unit cell, spaced by more than 51 nm, as would be required for full misfit strain relaxation. HRTEM and selected area ED characterization of the (Sr,La)Ox dots did not allow us to identify their structure with any of the Sr-La oxide phases reported in the literature [3], but appear to sustain a well defined three dimensional epitaxial relationship with the LSMO film. The matching distances provided by the nanodots to the LSMO film are d//=4.055 Å and d⊥=4.048 Å, parallel and perpendicular to the interface, respectively, and thus the misfit strains with the LSMO films are ε//=4.7% and ε⊥=4.5%. As a consequence the c parameter of the LSMO film becomes tensile stretched below the
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nanodots by ~+0.83%, resulting in an expanded LSMO unit cell volume ΔV/V∼+2.5%, as compared to the bulk LSMO unit cell. (Sr,La)Ox
Polytypoid phase LSMO
STO(001)
Figure 1. Low magnification TEM image showing the three phases system in the final microstructure. To get insights into the micro-structural trajectory of the film we have observed a sample quenched prior to the formation of the island. As shown in Figure 2, this sample is characterized by a high surface roughness and the coexistence of orthorhombic and rhomobohedral domains. The decrease of surface roughness of the LSMO film as the system evolves to the final microstructure suggests stress relief as the driving force for (Sr,La)Ox formation.
Figure 2. HRTEM image of a quenched microstructure prior to island formation. Local FFT patterns in A region and B region display the orthorhombic and rhombohedral variants respectively. 1. 2. 3.
C. Moreno, P. Abellan, A. Hassini, A. Ruyter, F. Sandiumenge, T. Puig, X. Obradors, In preparation. ICDD Diffraction Databases 2007. Newtown Square: International Centre for Diffraction Data. This work has been supported by EU (HIPERCHEM project), Generalitat de Catalunya (2005-SGR-0029 and CeRMAE) and by Spanish MEC (MAT2005-02047, FPU program)
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Microstructure of epitaxially strained LaCoO3 thin films L. Dieterle1, D. Gerthsen1 and D. Fuchs2 1. Laboratorium für Elektronenmikroskopie, Universität Karlsruhe, 76128 Karlsruhe, Germany 2. Institut für Festkörperphysik, Forschungszentrum Karlsruhe, 76021 Karlsruhe, Germany [email protected] Keywords: nano-twinning, LaCoO3, perovskite
Epitaxially strained LaCoO3 (LCO) films exhibit ferromagnetic ordering below a Curie temperature TC of 85 K in contrast to polycrystalline unstrained films which are paramagnetic at this temperature [1]. To correlate the observed magnetic properties with the microstructure, LCO films with different thicknesses of 10, 20, 50 and 100 nm were epitaxially grown by pulsed laser deposition. [001]-oriented (LaAlO3)0.3(Sr2AlTaO6)0.7 (LSAT) was used as substrate. Both, the substrates and the LCO, exhibit a perovskitetype structure. LCO is rhombohedrally distorted (aLCO = 3.80 Å, αLCO = 60.23 °, as bulk material) while LSAT is cubic (aLSAT = 3.86 Å). The lattice mismatch of ~ 1.5 % induces tensile strain in the LCO film. The microstructure of the samples was studied by transmission electron microscopy (TEM) combined with selected-area electron diffraction (SAED) and high-resolution TEM of plan-view and cross-section samples. Within the LCO thin films, strain is not reduced by the formation of interfacial dislocation but by nano-twinning along the [100] and [010] directions. Since both directions are equivalent on the surface of the cubic substrate, domains with nano-twins along both directions are observed within the thin films (cf. Figure 1). The domain structure is already formed at a LCO thickness of 10 nm. Each domain consists of twins with a width of about 15 nm (Fig. 1b, twin boundaries marked by arrows) and exhibit a tilt of ~1° towards the substrate (Fig. 2a). The layers are unstrained along the growth direction. However, the lattice is strained within the {110} in-plane directions and relaxation does not occur up to a layer thickness of 100 nm. The formation of this domain structure may be influenced by the ordering of the substrate [2], which was observed by SAED. Antiphase boundaries separate ordered regions (Fig. 2b) with doubled lattice constants, aLSAT, ordered = 7.72 Å. Size and shape of these ordered regions with partially nonconvex (curved out) contour-lines are similar to “size and shape” of the domains of the LCO thin films. 1. 2.
D. Fuchs, C. Pinta, T. Schwarz et al., Physical Review B75 (2007), 14402. H. Li, L. Salamanca-Riba, R. Ramesh et al., Journal of Materials Research 18(7) 2003, p. 1698
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Figure 1. a) Plan-view TEM bright-field image of the 10 nm LCO thin film taken with G g = (220) with domains oriented along the [100] and [010] directions and b) TEM bright-field image of the 100 nm LCO thin film in cross-section.
Figure 2. a) HRTEM image of the 20 nm LCO thin film along the [100] zone axis and b) TEM dark-field image of the 100 nm LCO layer taken with the superstructure reflection (1-11) along the [110] zone axis.
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Are the samples really flat? Influence of the supporting membrane on the magnetization of patterned micromagnets C. Dietrich and J. Zweck Institut für Experimentelle und Angewandte Physik, Universität Regensburg, Universitätsstr. 31, D-93053 Regensburg, Germany [email protected] Keywords: Lorentz microscopy, magnetic imaging, nanostructures
Lorentz transmission electron microscopy (LTEM) is a powerful tool to investigate the magnetization patterns of micro- and nanomagnets with high resolution. In the TEM the electron beam has to transmit the sample. Therefore, the specimens are usually grown onto thin Si3N4 membranes. AFM measurements show that the specimens are not perfectly flat but curved in a convex shape. The shape is in a good approximation paraboloidal with the center of a 1 µm x 1 µm Permalloy (Ni81Fe19) square only about 15 nm below the edges. Figure 1 shows an AFM image of elements grown onto bulk substrate and onto the free standing membrane. We assume mechanical tension due to different thermal expansion coefficients to be the origin of the curvature [1]. The ground state at remanence of square thin-film elements is typically the Kittel flux-closure domain structure, with four triangular domains and a Bloch line with perpendicular magnetization in the center. The domains are separated by 90° domainwalls aligned directly along the diagonals of the square. When a significant outof-plane magnetic field (0.1 T to 0.5 T) is applied, the magnetization patterns of the curved elements change remarkably. The domain walls get bent, so that the pattern looks “propeller-like”. Micromagnetic simulations based on the finite-element code TetraMag (which was also used in Ref. [2]), where the paraboloidal shape can precisely be modelled, show a perfect agreement with the experiments [1]. The magnetization patterns of thin Permalloy cylinders with a slight surface curvature are also modified in the presence of an out-of-plane field. For flat specimens the vortex is shifted perpendicular to an applied in-plane field, whereas for curved specimens the direction depends on the magnitude of an additionally applied out-ofplane field B (Figure 2, left). This effect can be qualitatively explained and reproduced by micromagnetic simulations. B can be locally split into a perpendicular component and a component BIP parallel to the specimen’s plane. Due to the slight curvature of the magnetic element the in-plane component forms a radial field with increasing magnitude from the center outwards (Figure 2, right), which modifies the applied homogeneous in-plane field and yields to the observed results. The out-of-plane field dependency gives us a further opportunity to modify the switching behaviour of nanodisks containing two or more antidots, which can be used as programmable logic or MRAM elements [3].
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 595–596, DOI: 10.1007/978-3-540-85226-1_298, © Springer-Verlag Berlin Heidelberg 2008
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1. 2. 3.
C. Dietrich, R. Hertel, M. Huber, D. Weiss, R. Schäfer, and J. Zweck, Phys. Rev. B. (2008), submitted R. Hertel, J. Appl. Phys. 90 (2001), 5752 M. Rahm, J. Stahl, and D. Weiss, Appl. Phys. Lett. 87 (2005), 182107
Figure 1. (a) AFM image showing Permalloy rectangles grown onto bulk Si (white) and on free standing Si3N4 membrane (dark). (b) Linescans from image (a): The bulk rectangles are uniformly flat but the elements grown on the membrane are curved down. (c) Scheme of the samples.
Figure 2. Lorentz images show a linear dependency of the vortex displacement angle on the applied out-of-plane field B (left) due to its inhomogeneous field BIP of in-plane components (right).
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HRTEM characterization of core-shell Fe@C and Fe@SiO2 magnetic nanoparticles prepared by the arc-discharge plasma method Rodrigo Fernández-Pacheco1, Manuel Arruebo1, Jordi Arbiol3, Clara Marquina2, Jesús Santamaría1, M. Ricardo Ibarra1,2 1. Nanoscience Institute of Aragon (INA) 2. Materials Science Institute of Aragon (CSIC-Zaragoza University), University of Zaragoza, 50009 Zaragoza, (Spain), 3. TEM-MAT, Serveis Cientificotecnics, Universitat de Barcelona, 08028 Barcelona, (Spain). [email protected]; [email protected] Keywords: Core-Shell, magnetic nanoparticles, HRTEM, EFTEM
The coating of functional magnetic nanoparticles with porous inorganic shells is a very important issue for various applications. Here we describe iron encapsulated in silica and carbon cages. The inorganic shell provides relevant properties: prevents aggregation and degradation of the metallic core, it can be functionalized to allow the binding of drugs or biomolecules to the carrier, is biocompatible, non-toxic, and possesses hydroxyl (i.e., silica)1 or carboxyl (i.e., carbon after chemical oxidation)2 surface groups. Those core-shell nanoparticles are good candidates for drug delivery applications; the drug release can be triggered by means of slow diffusion of the drug previously adsorbed on the porous matrix2,3. The arc-discharge method has been used as an innovative method to produce coreshell magnetic nanoparticles. The detailed structural and morphological characterization of our samples was carried out by means of transmission electron microscopy (TEM) and the power spectra obtained by using Fast Fourier Transform (FFT) algorithms on the high-resolution TEM (HRTEM) micrographs. In order to obtain the high-resolution TEM (HRTEM) results we used a field emission gun microscope Jeol 2010F, which works at 200 kV and has a point-to-point resolution of 0.19 nm. Electron energy loss (EELS) spectra were obtained in a Gatan Image Filter (GIF 2000) coupled to the Jeol 2010F microscope. Spectra were obtained with an energy resolution of 1.2 eV. 1. 2. 3.
R. Fernández-Pacheco, M. Arruebo, C. Marquina, M. R. Ibarra, J. Arbiol, J. Santamaría. Nanotechnology, Volume 17, pp 1188-1192. 2006 R. Fernández-Pacheco, C. Marquina, J. G. Valdivia, M. Gutierrez, M.S. Romero, R. Cornudellla, A. Labordad, A. Vitoria, T. Higuera, J.A. García de Jalón, M.R. Ibarra. Journal of Magnetism and Magnetic Materials. Volumen 311, pp 318-322. 2007. M. Arruebo, R. Fernandez-Pacheco, MR. Ibarra, J. Santamaría. anotoday. Volumen 2(3), pp 22-32. 2007.
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Si
Fe
Fe Fe Si
(100)Fe (022)Fe2O3
(02-3) Fe2O3 EFTEM SiO2 plasmon
Fe2O3 Hematite
Fe fcc
Figure 1. EFTEM and HRTEM patterns for the Fe-silica nanoparticles. Blue: SiO2 plasmon. Red: Fe fcc HRTEM structure. Green: Fe2O3 hematite HRTEM structure (see Ref. 1).
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Nanofabrication of ferromagnetic nanotips and nanobridges by 2D and 3D electron-beam cutting T. Gnanavel, Z. Saghi, Y. Peng, B.J. Inkson, M.R.J. Gibbs and G. Möbus Dept. of Eng. Materials, University of Sheffield, Mappin Street, Sheffield, S1 3JD, UK. [email protected] Keywords: hole drilling, nanofabrication, electron tomography, nanobridges, nanotips
Electron beam nanofabrication is a powerful technique to generate finest structures by atom ablation while concurrently observing the progress by TEM imaging [1-3]. We first explore fabrication of nickel nanobridges starting from a cross-sectional TEM specimen of a Ni thin film on Si substrate using several focused electron beam milling steps on a JEM 3010 LaB6 microscope at 300kV. The purpose of such devices is for magnetotransport studies. In Figure 1, progress towards the fabrication of a bridge and finally breakdown of the bridge is sampled. In extending such thin film plan view studies, we develop e-beam nanofabrication into a 3D STEM nanofabrication technique with special application to cone shaped metal nanotips. The nanowire/tip axis is mounted along the goniometer axis and perpendicular to the electron beam. In a single viewing direction, repeated e-beam cutting could only turn the conical tip into a disk shaped tip. By using a high-tilt tomographic specimen holder, the cutting can be iterated in pairs of e.g. 90O oriented sample angles. Most importantly, this 3D nanofabrication can be immediately inspected by a tomographic imaging tilt series (after the beam is spread to stop further cutting). Experimental demonstrations are shown on a JEM 2010F FEG-TEM at 200kV using a home-made Ni tip, etched from a Ni wire. Cutting/drilling by electron beam is performed with the largest condenser aperture, spot size 1, and fully focused beam in TEM or CBED mode. The Ni-tip is mounted with full rotational symmetry [4] in a hightilt tomography holder [5]. Figure 2a shows the hole and line, drilled for test purpose (arrows 1 and 2). The radical reduction of tip radius for a nickel tip can be seen in Figure 2b following a first cut at the centre of tip. Examples of cutting of holes through the middle as well as cuts into the side of the wire are collected in Figure 2c. After rotation of the specimen by 90 degrees, the cuts and holes can be analysed (Figure 2d) and further holes (Figure 2e) perpendicular to the first series can be drilled. Holes drilled before and observed in perpendicular direction appear as bands of reduced thickness, sometimes including redeposition of Ni outside the tip (Figure 2d&2e). [6] 1. 2. 3. 4.
K. Furuya, K. Mitsuishi, M. Shimojo and M. Takeguchi., Rev. Adv. Mater. Sci. 5 (2003) p. 381. H. W. Zandbergen, R. J. H. A. van Duuren, P. F. A. Alkemade, G. Lientschnig, O. Vasquez, C. Dekker and F. D. Tichelaar, Nano Lett., 5, (2005) p. 549. M. D. Fischbein and M. Drndic, Nano Lett., 7, (2007) p. 1329. X. Xu, Z. Saghi, G. Yang, Y. Peng, B.J. Inkson, R. Gay and G. Möbus, Mater. Res. Soc. Symp. Proc., 928E (2007) 0982-KK02-04.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 599–600, DOI: 10.1007/978-3-540-85226-1_300, © Springer-Verlag Berlin Heidelberg 2008
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G. Möbus, B.J. Inkson, I.M. Ross and R. Morrison, Microscopy & Microanalysis, 10 (Suppl 2) (2004) p. 1196. This work was supported by EPSRC under grant number GR/S85689/01.
(a)
(b)
(c)
(d)
Figure 1. In-situ observation of contact modification of nanoconstriction while irradiating by e-beam. (a) Drilled nanoconstriction. The width of the constriction is 15nm. Scale bar: 20nm. (b-c) Image captured a few seconds before rupture with two different magnifications. Scale bars: (b)-10nm and (c)-5nm. (d) Just after rupture. The width of the gap is 3nm (indicated by square box). Scale bar: 5nm. JEM 3010 LaB6, 300kV.
(a) 0O
(b) 0O
3
12 8
4
56
9 (c) +45O
7
(d) -45O
(e) -45O
10
Figure 2. (a) Ni nanotip before sharpening by focused electron beam (view at 0 degrees). An additional hole and line have been drilled for test purpose (arrows 1 and 2). (b) After sharpening (view at 0 degrees). Arrow 3 indicates material protruding outside original silhouette. (c) View of tip after rotation by +45O and drilling of additional holes through centre (arrow 4) and cuts through the side (arrows 5-9). (d) View after rotation by 90 degrees without further milling. Traces of hole 1 and extrusions (arrows). (e) Drilling of a further hole with orientation perpendicular to hole 1 (arrow 10). JEM 2010F FEGTEM, 200kV.
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An investigation into the crystallization of the MgO barrier layer of a magnetic tunnel junction V. Harnchana1, A.P. Brown1, R.M. Brydson1, A.T. Hindmarch2, and C.H. Marrows2 1. Institute for Materials Research, University of Leeds, UK 2. School of Physics and Astronomy, University of Leeds, UK [email protected] Keywords: magnetic tunnel junction, CoFeB/MgO/CoFeB, crystallization
Magnetic tunnel junctions (MTJs) are multilayer structures, consisting principally of two ferromagnetic layers separated by an insulating layer, that exhibit a tunnelling magnetoresistance (TMR) phenomena. MTJs with a large TMR value are general requirements for magnetic random access memory (MRAM) applications and the next generation magnetic sensors. Coherent lattice matching at the interfaces of a CoFe/MgO based MTJ (CoFe (100)[110]/MgO(100)[100]/CoFe(100)[110]) is predicted to give a very high TMR value of over 6000%[1]. Sputter-deposited CoFeB/MgO MTJ devices exhibiting lattice matching between the CoFeB and MgO that have a large TMR ratio of 500 % have been reported recently [2]. The growth and crystallization mechanism of this active trilayer plus the relationship of the sputtering parameters with the film structure have not been thoroughly studied. TEM characterization of the device is one way to explore this. TEM of MTJs is typically performed on a cross-section structure, the preparation of which is time-consuming and the analysis limited to only a specific area of the junction. Here we present a fundamental study of directly sputter-deposited thin films onto a TEM support grid (using the combination of dc-rf magnetron sputtering). In order to verify this planar analysis technique is representative, a TEM cross-section specimen was prepared from a simple MTJ deposited onto a Si wafer substrate for comparison (Figure 1). Both a single CoFeB layer and a trilayer of CoFeB/MgO/CoFeB were produced as planar TEM specimens (Figure 2) using the same sputter-deposition conditions under which the MTJ device was prepared. TEM imaging, selected area electron diffraction (SAED) and energy dispersive X-ray (EDX) spectroscopy, electron energy loss spectroscopy (EELS), and energy filtered TEM (EFTEM) techniques were employed to study the nanostructure. TEM analysis of the active trilayer, CoFeB/ MgO/CoFeB from the cross-section shows that a discontinuous layer of MgO crystallites forms in the device (Figure 1). Furthermore, the MgO crystallites do not lie in (100) orientation with respect to the interface plane, although where the (200) planes are visible the lattice fringes do apparently extend into the upper CoFeB layer (central crystallite in Figure 1). It is clear from a TEM image of the active trilayer planar specimen (Figure 2) and the associated diffraction pattern that the sputter-deposited MgO layer forms isolated crystallites with random orientations in an amorphous (CoFeB and carbon support film) matrix (elemental mapping confirms that the MgO layer is continuous). Investigation of the optimum route for deposition of a textured MgO (100) layer will be undertaken by
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 601–602, DOI: 10.1007/978-3-540-85226-1_301, © Springer-Verlag Berlin Heidelberg 2008
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systematically varying the sputtering conditions. How that layer induces the preferredorientation crystallization of the surrounding CoFeB layer during annealing will then be considered. 1. 2.
X G Zhang and W H Butler, Phys. Rev. B 70(2004), p. 172407 Y M Lee, et al, Appl. Phys. Lett 90 (2007), p. 212507.
MgO, d(111) = 2.43 Å d(200) = 2.18 Å
d(111) = 2.43 Å
CoFeB/MgO/CoFeB
Figure 1. High magnification TEM image of a cross-section of the MTJ produced by the School of Physics, Univ. of Leeds. The image shows the active trilayer region of CoFeB[3-4nm]/MgO[7-8nm]/CoFeB[3-4nm]; the crystallites are identified as MgO with the (200) and (111) planes visible. The interfaces of CoFeB are relatively rough and no clear CoFeB/ MgO boundaries are observed. a)
b)
MgO (111) (200) (311) (400) (420)
Figure 2. a) TEM image of a planar sample of the CoFeB/MgO/CoFeB trilayer sputtered deposited under the same conditions as the MTJ in Figure 1. In a) isolated crystalline regions are visible and the SAED pattern of the film b) shows that the crystallites can be indexed to MgO with a random orientation.
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FeCoAlN films with induced magnetic anisotropy A. Lančok1,2, M. Klementová1, M. Miglierini1,3, F. Fendrych2, K. Postava4, J. Kohout5, and O. Životský4 1. Institute of Inorganic Chemistry of the ASCR, v.v.i., 250 68 Husinec-Řež 1001, CZ 2. Institute of Physics of the ASCR, v.v.i., Na Slovance 2, 18 221 Prague 8, CZ 3. Dept. of Nuclear Physics and Technology of STU, Ilkovičova 3, 812 19 Bratislava 4. Dept. of Physics, Technical University of Ostrava, 17.listopadu 15, 708 33 Ostrava 5. Faculty of Math and Physics, Charles University, V Holešovičkách 2, Prague, CZ [email protected] Keywords: thin films, nanoparticles, plasma-jet deposition, Mössbauer spectroscopy, NMR, TEM
Nowadays, an intensive research of new magnetic materials with high frequency permeability has been stimulated by the demand for electronic devices, such as microinductors, transformers and magnetic recording heads, operating in the GHzfrequency range [1]. Soft magnetic nanogranular films consisting of magnetic grains embedded in an insulating matrix are appropriate candidates for these applications due to: (a) the intrinsic maze structure of nanogranular films leads to high resistivity which reduces eddy current losses, (b) the ferromagnetic resonance (FMR) can be shifted to higher frequencies in materials with high saturation magnetization and high anisotropy field. Research on nanogranular ferromagnetic FeCoAlN films is presented. The FeCoAlN films with induced magnetic anisotropy were prepared by a specially designed UHV plasma-jet system with DC hollow-cathode discharge [2]. Plasma deposition process was performed by reactive sputtering of combined Fe50Co50+Al nozzle in an Ar+N2 gas mixture flow on water-cooled Si, SiO2/Si and glass substrates. TEM, Mössbauer spectroscopy, magneto-optical Kerr effect (MOKE) measurements, and NMR were used to characterize the samples in detail and to study the effects of the conditions of synthesis on sample properties. The films have thickness of about 600 nm. They are composed of nanoparticles of cubic FeCo, 5-20 nm in size, embedded in amorphous matrix containing Al, N and O (Figure 1). According to electron diffraction, an additional crystalline phase is present, most likely cubic Fe5Co3. According to MOKE, in-plane hysteresis loops of as-deposited samples in the longitudinal and transversal direction to the applied field show excellent soft magnetic properties with the coercive field of about 4 Oe and in-plane magnetic uniaxial anisotropy. The coherent rotation of magnetization followed by nucleation of domains magnetized along the easy axis of the film is observed [3]. The conversion electron Mössbauer spectra (CEMS) of FeCoAlN film were decomposed into 3 sextets with hyperfine fields (Bhf) of about 32, 33.8, and 35 T and one doublet. The spectra indicate that the orientation of local magnetic moments on the surface is random for the sample with 1:1 Fe/Co ratio, while the magnetization lies in the plane of the layer for samples with Fe/Co smaller than 1 (0.91-0.95).
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According to NMR spectra of 57Fe and 59Co nuclei, the increase of average Bhf at Co by ~3.5T with respect to Co powder agrees with the effect of Fe atoms as the nearest neighbours of resonating 59Co nuclei. Similar effect is responsible for the increase of Bhf at 57Fe nuclei due to the Co nearest neighbours. The average Bhf from NMR for 57Fe and 59Co in the films are higher than for the powders of pure elements and their distribution is rather broad which indicates the presence of an amorphous phase and/or very small crystalline particles as confirmed by HRTEM. 59
1. 2. 3. 4.
K. Seemann, H. Leiste and V. Bekker, J. Magn. Magn. Mater. 302 (2006), p. 321. F. Fendrych, L. Kraus, O. Chayka, P. Lobotka, I. Vávra, J. Touš, V. Studnička, and Z. Frait, Monatshefte J. Chemie 133 (2002), p. 773. O. Životský, F. Fendrych, L. Kraus, K. Postava, O. Chayka, L. Halagačka and J. Pištora, J. Magn. Mater. 316(2) (2007), p. e858. We kindly acknowledge the support of Grant Agency of the ASCR (projects KAN 400100653 and IAA200100701), and Ministry of Education of the Czech Republic (projects VZ 0021620834 and OC08030 (COST P17)).
Figure 1. TEM micrographs of the FeCoAlN film. Typical microstructure of the film consists of FeCo nanoparticles (electron diffraction pattern as an inset) and amorphous matrix containing Al, N, O (on the left). On the right, high resolution micrograph of cubic FeCo nanoparticle along [111].
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The martensitic microstructure of 5M and NM martensites in off-stoichiometric Ni2MnGa ferromagnetic shape memory alloys Pallavi Sontakke, Amita Gupta and Madangopal Krishnan Materials Science Division, Bhabha Atomic Research Centre, Mumbai-400085, India. [email protected] Keywords: martensite, shape memory, Ni-Mn-Ga
In recent years, extensive studies have been performed on many aspects of the NiMnGa alloys, such as crystal structure, phase transformation, magnetic field induced strain, magnetic properties, mechanical behaviour and the effects of magnetic fields on martensitic transformation. Depending upon the composition, the most important martensites that exist in off-stoichiometric Ni2MnGa alloys are the modulated 5M and 7M, and, the non-modulated tetragonal NM ones. In this study, a systematic investigation was carried out on the 5M and NM martensite in well characterised Ni50Mn28.9Ga21.1 and Ni53Mn25Ga22 alloys, respectively using transmission electron microscopy (TEM) and X-ray diffractometry for determining crystal structure, substructure, intervariant interfaces and self-accommodating microstructure. It was established that both NM and 5M are internally twinned martensites with tetragonal and monoclinic crystal structures, respectively, Figures1-4. The crystallographic features of the martensitic microstructures as observed by TEM are found to be in agreement with those predicted by computations based on the phenomenological theory of martensite crystallography.
Figure 1. [010] SADP shows five layered modulation in [001] direction .
Figure 2. Internally twinned 5M martensitic plates with coherent intervariant interfaces.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 605–606, DOI: 10.1007/978-3-540-85226-1_303, © Springer-Verlag Berlin Heidelberg 2008
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Figure 3. Twinned diffraction pattern obtained from the internal twins.
Figure 4. Twin related variants of NM martensite plates.
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TEM characterization of nanometer-sized Fe/MgO granular multilayer thin films grown by pulsed laser deposition C. Magén1, E. Snoeck1, A. García-García2, J.A. Pardo3,4, P.A. Algarabel2, P. Štrichovanec2, A. Vovk3, L. Morellón3, J.M. De Teresa2 and M.R. Ibarra3 1. Centre d’Elaboration de Matériaux et d’Etudes Structurales, CNRS, 31055-Toulouse, France 2. Instituto de Ciencia de Materiales de Aragón, Universidad de Zaragoza-CSIC, 50009Zaragoza, Spain 3. Instituto de Nanociencia de Aragón, Universidad de Zaragoza, 50009-Zaragoza, Spain 4. Dpto. de Ciencia y Tecnología de Materiales y Fluidos, Universidad de Zaragoza, 50018-Zaragoza, Spain [email protected] Keywords: transmission electron microscopy, granular materials, pulsed laser deposition
Research of granular thin films composed by magnetic single-domain nanoparticles dispersed in a non-magnetic matrix is currently very active because their magnetic and magnetotransport properties suggest various technological applications. Discontinuous metal insulator multilayers (DMIMs) belong to this class of materials. DMIMs consist of metallic layers with different degrees of discontinuity intercalated between insulating spacer layers. When the metallic layer consists of closely spaced ferromagnetic grains, different types of collective behaviour may appear due to magnetic dipolar or exchange interactions [1, 2]. In fact, the magnetic behaviour can be controlled by modifying the nominal thickness of the deposited ferromagnetic metal [3]. Therefore, microscopic characterization is essential in tailoring the magnetic properties of these systems. In this work we have carried out a transmission electron microscopy (TEM) characterization of granular Fe/MgO multilayers. The samples have been grown by sequential pulsed laser deposition (PLD) of Fe and MgO films on Corning glass substrates. Specimens with Fe nominal thickness in the range of tFe=0.4-1.25 nm have been prepared, but the MgO thickness of tMgO=3 nm has been fixed. Phase contrast and Z-contrast TEM images have been obtained using a Cs-corrected Tecnai F20 and a Tecnai F30 microscope [4], respectively, on cross section specimens to characterize the layered structure of the stacking (see Fig.1). In specimens with tFe=0.4 nm and 0.8 nm, bright field and dark field images have evidenced the polycrystalline nature of the Fe layers and the MgO spacers. In the case of MgO, a tendency to columnar growth through the Fe layers has been observed. The discontinuous nature of the Fe layers is not clear from these images due to the nanometric size of the grains, which is likely to be smaller than the thickness of the areas of interest in the cross section TEM samples.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 607–608, DOI: 10.1007/978-3-540-85226-1_304, © Springer-Verlag Berlin Heidelberg 2008
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To clarify this point, plane view bilayers of Fe/MgO with the same nominal Fe thickness have been deposited on carbon grids and observed in a Philips CM30 microscope. The specimen with tFe=0.8 nm is shown in Fig. 2(a). In this case MgO grows in amorphous phase, whereas discrete Fe nanometer-sized grains are easily distinguishable. The size distribution and dispersion of Fe grains have been measured and fitted to log-normal function with a maximum at 4.5 ± 0.3 nm, see Fig. 2(b). Results for different Fe thickness will be discussed in correlation with magnetic measurements. 1. 2. 3. 4. 5.
Present address: Oak Ridge National Laboratory, P.O. Box 2008, Oak Ridge, TN 37831 US. Xi Chen et al., Phys. Rev. B 72, 214436 (2005). S. Bedanta et al., Phys. Rev. Lett. 98, 176601 (2007). W. Kleemann et al., Phys. Rev. B 63, 134423 (2001) We kindly acknowledge Dr. Dong Tang for his assistance in the Tecnai F30 experiments.
(a)
(b)
Figure 1. (a) Cross section HRTEM image of the Fe/MgO multilayer with nominal thicknesses tFe=0.81 nm. (b) Z-contrast image of the same specimen.
(a)
(b)
Figure 2. (a) Plane view bright field image of the Fe/MgO bilayer with nominal thickness tFe = 0.81 nm. (b) Fe grain size distribution showing the best log-normal fit.
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Structural modification and self-assembly of nanoscale magnetite synthesised in the presence of an anionic surfactant S. Malik1, I.J. Hewitt2, A.K. Powell1,2 1. Institute of Nanotechnology, Forschungszentrum Karlsruhe GmbH, PO Box 3640, D-76021 Leopoldshafen-Eggenstein, Germany. 2. Institut für Anorganische Chemie,Universität Kalrsruhe, Engesserstrasse 15, D-76131 Karlsruhe, Germany. [email protected] Keywords: magnetic nanoparticles, catalyst
Many biological and industrial processes are crucially dependent upon the absorption of surfactants from an aqueous phase onto a solid surface [1,2]. At the heart of this physical chemical process is the alteration of the interface properties caused by the adhesion and aggregation of the surfactant molecules at the solid surface. The electrostatic interaction between a charged surfactant head-group and the intrinsic charge that many solids have when they are immersed in an aqueous phase, often derives from the attraction of aqueous surfactants to a solid phase. Thus, the driving force for the initiation of monomer absorption and subsequent aggregation of the surfactant molecule through dispersion and van der Waals forces is the result of the electric potential in the interface region. These interactions can result in dramatic changes in the properties of the solid surface and the interfacial region as the surface charge is altered [3]. Synthesis of magnetite (Fe3O4) in the presence of the surfactant sodium dodecyl sulphate (SDS) gives rise to a variety of nanoscale morphologies, some of which look remarkably similar to magnetite found in organisms, suggesting that similar processes may be involved. So, taking our inspiration from biology, where templates produce magnetite of defined shapes and sizes, we have been interested in investigating how surfactant molecules can similarly influence nanoscale magnetite formation. The structurally modified nanoscale magnetite were characterised by Scanning Electron Microscopy (SEM). In this paper, we report structural modification and self-assembly of magnetite in the presence of the anionic surfactant SDS. SDS is commonly used to mimic hydrophobic binding environments such as cell membranes [4] and has recently been used to study the folding and thermal stability of cytochrome c (cyt c) a biologically important electron transfer system [5]. 1. 2.
A. W. Adamson, Physical Chemistry of Surfaces, 5th ed.; John Wiley & Sons: New York, 1990. D. Myers, Surfaces, Interfaces, and Colloids. Principles and Applications, 2nd ed; John Wiley & Sons: New York, 1999.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 609–610, DOI: 10.1007/978-3-540-85226-1_305, © Springer-Verlag Berlin Heidelberg 2008
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K. A. Becraft, F. G. Moore, G. L. Richmond, J. Phys Chem B 2003, 107, 3675-3678. M. N. Jones, Biological interfaces: An Introduction to the Surface and Colloid Science of Biochemical and Biological System. Elsevier, Amsterdam 1975. Q. Xu, T. A. Keiderling, Protein Science 2004,13, 2949–2959.
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Electron microscopy phase retrieval of perpendicular magnetic anisotropy (PMA) FePd alloys A. Masseboeuf1, C. Gatel2, A. Marty1, E. Snoek2 and P. Bayle-Guillemaud1 1. INAC/SP2M, Cea Grenoble, 17 rue des Martyrs, 30854 Grenoble Cedex 9 - FR 2. CEMES, CNRS, 29 rue Jeanne Marvig, 31055 Toulouse Cedex - FR [email protected] Keywords: Electron holography, Transport of intensity equation, Perpendicular magnetic anisotropy
Among new materials designed for magnetic logic or future MRAM devices, the chemically ordered, self organized, FePd alloys are of the best candidates. We will present analysis of the magnetic configurations of these alloys performed by Electron holography [1] and Lorentz microscopy [2]. Comparison of different ways for imaging magnetic distribution in a TEM will also be discussed. Magnetic FePd (L10)/FePd (disorderd) bi-layers have been investigated by Electron Holography both in plane and cross-sectional geometry. This alloy epitaxially grown on MgO(001) substrate in its ordered L10 phase exhibits a magneto-crystalline uniaxial anisotropy. The magnetic anisotropy, resulting from such a chemical organisation, gives rise to “up” and “down” domains configuration. The vanishing anisotropy of the second FePd disordered layer induces an alignment of the domains in a stripe configuration after magnetic saturation. A major problem to image magnetic domains in a TEM by measuring the phase shift of the e-beam is that the magnetic information is mixed with the electrical one. As described by Aharonov and Bohm [3], the phase shift of the electron wave crossing a sample is:
φ = C E ∫ Vint dz +
e Az dz =∫
where CE is a electron-energy related constant, Vint represents the mean inner potential of the material (the electrostatic component) and Az is the z-component of the vector potential, resulting of the magnetic induction. The aim of a phase retrieval process in electron microscopy is to separate the two contributions. In this work, both electron holography and Fresnel focal series will be presented with particular way of removing the electrical contribution. For the in-plane geometry (Figure 1) the magnetic induction mapping is obtained by subtracting a vectorial constant due to a gradient thickness of the sample [4]. The magnetic configuration of the bi-layers studied in cross-sectional sample will be resolved by electron holography (Figure 2). Here the electrical and magnetic contributions to the phase shift will be separated from calculations on the two holograms obtained by flipping the sample out of the microscope. The “up” and “down” domain configuration is clearly observed with flux closure within the FePd disordered layer. Micromagnetic configuration performed to simulate such configuration will be presented.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 611–612, DOI: 10.1007/978-3-540-85226-1_306, © Springer-Verlag Berlin Heidelberg 2008
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1. 2. 3. 4.
D. Gabord, Nature 161 (1948) p. 777–778 J.N. Chapman, Journal of Physics D : Applied Physics 17 (1984) p. 623-647 Y. Aharonov, and D. Bohm, Physical Review 115 (1959) p. 485–491 A. Masseboeuf, et al. submitted to Journal of Applied Physics
Figure 1. Fresnel image under-focused (left) and mapping of the magnetic induction distribution (right) of FePd(ord.)/FePd(disord.) alloy exhibiting a stripe configuration.
Figure 2. Up: Magnetic configuration to the phase shift. Down: electrostatic contribution. Iso-phase lines are displayed as 0.25-rad for the magnetic part and 1-rad for the electrical part.
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Magnetic domain wall propagation in nanostructures of alloys with perpendicular magnetic anisotropy A. Masseboeuf1, A. Mihaï2, J.P. Attané2, J.C. Pillet2, P. Warin2, A.L. Vila2, G. Gaudin3, M. Miron3, B. Rodmacq3, E. Gautier3, A. Marty2 and P. Bayle-Guillemaud1 1. INAC/SP2M/LEMMA, Cea/UJF Grenoble, 17 rue des Martyrs, 30854 Grenoble 2. INAC/SP2M/NM, Cea/UJF Grenoble, 17 rue des Martyrs, 30854 Grenoble Cedex 9 3. INAC/SPINTEC, Cea Grenoble, 17 rue des Martyrs, 30854 Grenoble Cedex 9 [email protected] Keywords: Domain wall propagation, Perpendicular Magnetic Anisotropy
Among the new challenges of the domain walls (DW) propagation, the control of DW movements induced by spin injection is extensively studied. Magnetic Transmission Electron Microscopy is a powerful tool for imaging the magnetic domain and analyse in situ their movement. It offers a nanoscale resolution and is very sensitive to magnetic induction inside the sample. Most of the investigations on DW propagation using Lorentz Microscopy (LTEM) are done on planar configuration [1-2]. Recently, the first observations of current-excited DW propagation, the well known Spin-Torque effect, were achieved [2]. The work presented in this paper concerns an original approach of the study of DW propagation in CoPt multilayers with perpendicular magnetic anisotropy. The Lorentz Microscopy is used to reveal domain walls by tilting the sample, giving rise to in-plane magnetic components, essential for magnetic contrast. The multi-layers were sputtered on silicon nitride windows. They consist of a stack of [Co/Pt]n multilayer (the Co thickness being <1 nm to keep the perpendicular anisotropy). .The number of layers n in the stacking determines the total Co volume in the sample. This value is an important parameter for the contrast in the LTEM images and impacts the Signal-To-Noise Ratio. Indeed, for low value of n, high defocus is needed to get Fresnel contrast on the domain walls. Moreover, for thick film the magnetic distribution inside the layers becomes complex. Figure 1 presents a magnetization process of such a film observed by Lorentz microscopy with a huge défocalisation (more than 400 µm). The foils have then been patterned by Focused Ion Beam (FIB) to define straight lines and dedicated structures where domain walls are supposed to trap and/or to expand. Irradiation inside the patterns can be found near etched areas as described in [3]. We will show how that irradiation process is suitable for injecting domain walls in the nanostructures. Figure 2 presents a typical pattern etched for DW propagation. Two pads are separated by a 500 nm line. The DW are easily nucleated inside the pads of the structure by a perpendicular applied field and can then be injected in the lines. 1. 2. 3.
D. McGrouther et al., Applied Physics Letters 91 (2007), p. 022506 Y. Togawa, et al., Applied Physics Letters 92 (2008), p. 012505 P. Warin, et al., Journal of Applied Physics 90 (2001), p. 3843
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 613–614, DOI: 10.1007/978-3-540-85226-1_307, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. Fresnel images of magnetization reversal of the Co/Pt multi-layer. The images are 10 x 10 µm. The sample has been tilted of 20° to give rise to Fresnel contrasts. The défocalisation used here is about 1 mm. The applied field corresponding to each image is indicated.
1µm
Figure 2. CoPt layer patterned using a Focused Ion Beam (MEB image in the inset). We can observe some domain walls in the pad on both sides. The deformation of the pattern on the LTEM image is due to the tilt of the sample.
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The effect of annealing in the microstructure and magnetic properties of NiCuZn ferrites D. Sakellari1, V. Tsakaloudi2, V. Zaspalis2 and E.K. Polychroniadis1 1. Solid State Physics Section, Department of Physics, Aristotle University of Thessaloniki-Greece 2. Laboratory of Inorganic Materials, Center for Research and Technology-Hellas, Thessaloniki-Greece [email protected] Keywords: ferrites, microstructure, stress, magnetic properties
NiCuZn ferrites are widely used in high frequency applications due to their high electrical resistivity and low losses, under low production costs [1,2]. As the properties of ferrite materials are influenced by their microstructure, which in turn is sensitive to the manufacturing procedure [3,4], the knowledge of the relation between process parameters and the microstructural characteristics is of great importance. In the present work, NiCuZn ferrite samples with the stoichiometry (Ni0,30 Cu0,07Zn0,63)Fe1,93O4 were synthesized by the mixed oxide process and sintered between 1025-1125°C. Each sintered sample was further annealed for 30min at 50°C below the applied sintering temperature. For all samples before and after annealing, a microstructural study was carried out using Transmission Electron Microscopy (TEM), while the magnetic properties were measured by a programmable unit with an impedance-gain analyzer. Concerning the samples before annealing, TEM reveals the existence of defects such as dislocations (Figure 1a) for samples with sintering temperatures 1025-1050°C, subgrain boundaries (Figure 1b) for all samples and phase separation, forming a compositional modulation-like contrast (Figure 1c). Dislocations, which are evidence of internal stress of the material, seem to disappear after the annealing. The subgrain boundaries, which indicate a first annealing during sintering, remain unaffected. Compositional modulations are also evident only for samples before the annealing and disappear completely after this procedure. The evaluation of the magnetic properties (Table 1) reveals that there is a significant decrease of the power losses for each sample after annealing, while the initial permeability levels remain unchanged. A mechanism through which the annealing process enhances the relaxation of the present stresses, being quantitatively evaluated by the reduction of the power losses, is proposed. In fact, there is a gradual reduction of the loss improvement due to the annealing, which becomes zero when the sintering conditions have been optimised. This mechanism of stress relaxation as well as the material homogenization are confirmed by the TEM observations. 1. 2.
Y.Matsuo, M. Inagaki, T. Tomozawa, F. Nakao, IEEE Trans. Magn. 37 (2001) 2359 I.Z. Rahman, T.T. Ahmed, J. Magn. Magn. Mater. 290-291 (2005) 1576-1579
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 615–616, DOI: 10.1007/978-3-540-85226-1_308, © Springer-Verlag Berlin Heidelberg 2008
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3. 4. 5.
A.Verma, T.G.Goel, R.G.Mendiratta and M.I.Alam, Mat. Sci. and Eng., B60 (1999) 156-162 M.H. Khedr, A.A.Omar, M.I. Nasr and E.K. Sedeek, J.Anal. Appl. Pyrol. 76 (2006) 203-208 This work was partially supported by the General Secretariat for Research and TechnologyPENED 2003
Figure 1. Bright Field TEM image with (a) dislocations, (b) subgrain boundary and (c) compositional modulation from sample sintered at 1050 °C (a,b) and 1125 °C (c) Table 1. Magnetic properties of samples before and after annealing Sintering temperature ( ◦C )
ui at 30oC, 10 kHz 0.1 mT
1025 1025 ann 1050 1050 ann 1075 1075 ann 1100 1100 ann 1125 1125 ann
493 498 772 759 1073 933 1047 1156 1175 1163
tand/ui at 30oC, 1 MHz, 0.1 .T 53 58 67 67 79 84 105 119 135 134
Pv at 25C, 50 kHz, 150 mT 1085 1058 646 624 466 447 392 388 362 362
% reduction
2.49 3.41 4.08 1.02 0.00
Bsat at 25C, 25 kHz, 250 A/m 202 208 250 250 265 275 276 277 277 277
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Microstructural and compositional analyses of nano-structured Co-Pt thin films Z. Samardžija, K. Žužek Rožman, S. Kobe Department for Nanostructured Materials, Jožef Stefan Institute, Jamova cesta 39, SI-1000 Ljubljana, Slovenia [email protected] Keywords: CoPt thin films, AFM, microanalysis, WDXS
Co-Pt alloys have attracted a lot of interest from the applied-magnetism community because of their versatile magnetic properties and good corrosion resistance, which fit in a wide range of potential applications, such as magneto-optical recording, perpendicular magneto-recording, hard magnetic components and, most recently, in MEMS. Co-Pt thin films can be prepared using the electro-deposition method from a single electrolyte [1]. Films with a composition close to Co50Pt50 were deposited potentiostatically on Auand Ag-coated glass substrates at applied voltages from -0.90 V to -1.15 V vs. Ag/AgCl electrode, for 20 min. The thickness, the surface morphology and the chemical composition of the deposited Co-Pt films were investigated using scanning electron microscopy (SEM), atomic force microscopy (AFM) and electron-probe microanalysis (EPMA) with energy-dispersive (EDXS) and wavelength-dispersive x-ray spectroscopy (WDXS). The characteristic morphology of a cross-section of Co-Pt films is shown in Figure 1a. The as-deposited Co-Pt films are nano-structured and consist of rounded grains with a size below 50 nm. The film thickness was proportional to applied electric potential and varied between 150 nm and 300 nm. The surfaces of all the films were visually smooth and mirror-like. A detailed AFM analysis of the surface morphology allowed us to determine the size and the shape of nano-grains with higher confidence, as shown in Figure 1b. All the films had grains in the range 20-50 nm. The smoothness of the film surfaces was consistent with an average roughness (Ra) of 2.4 ± 0.5 nm. For EPMA the experimental setup was optimized to obtain a reliable compositional analysis of the Co-Pt thin films using the conventional "bulk" EPMA approach. The results of Monte Carlo simulations for electron interaction volume and Φ(ρz) x-ray depth distributions showed that a reliable microanalysis can be performed by measuring the intensities of the low-energy Co-Lα and Pt-Mα spectral lines at low primary beam energies. Calculations for a 6-keV beam showed that the depth of x-ray generation in the Co-Pt material is about 100 nm; consequently, the whole x-ray excitation volume remains within the Co-Pt layer. Qualitative elemental analyses with EDXS confirmed the presence of Co and Pt (Figure 2a). Small amounts of Cu detected in a layer originate from the impurities in the electrolyte. The WDXS measurements for quantitative analysis were performed at 6-kV voltage, 90-nA beam current and a 40o take-off angle. The Pt-Mα line was analyzed on a PET crystal and the Co-Lα line on a TAP crystal. The standards were pure Pt and Co. The measured k-ratios were quantified using
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 617–618, DOI: 10.1007/978-3-540-85226-1_309, © Springer-Verlag Berlin Heidelberg 2008
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various Φ(ρz) matrix-correction methods: PROZA, PAP, XPP and XPHI. The counting times were set to ensure high analytical precision with a standard counting error of < 1 % relative. The concentrations of Co and Pt calculated using four matrix-correction programs were fully consistent and within a relative standard deviation of ± 1.9 %. The achieved non-normalized analytical totals fall within 98-99 %, which verifies the correctness of the applied analytical method and the high certainty of the quantitative results. The concentrations of Co and Pt vary between the samples, as shown in Figure 2b. The Co/Pt ratio is closest to the nominal one in samples #5 and #1, whereas in other samples it scatters from 0.76 to 1.35. The results indicate that the chemical composition of electro-deposited Co-Pt thin films is sensitive to slight changes in the applied voltage, when other process parameters were kept constant. Therefore, precise control of the voltage is required to get the desired Co-Pt alloy stoichiometry. The WDXS analysis is a reliable and suitable tool for accurate, quantitative, compositional analyses of thin Co-Pt films with thicknesses down to 100 nm. 1.
P.L.Cavallotti, M. Bestetti, S. Franz, Electrochimica Acta, 48 (2003), p. 3013-3020
b
a
Figure 1. (a) SEM micrograph of fractured Co-Pt thin film cross section, (b) AFM image of thin film surface morphology. a
b
Figure 2. (a) EDXS spectrum acquired from the Co-Pt film, (b) results of quantitative WDXS analysis for Co and Pt concentrations in analyzed Co-Pt thin films.
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L10-type ordered structure of FePd nanoparticles studied by high-resolution transmission electron microscopy K. Sato1, T.J. Konno1 and Y. Hirotsu2 1. Institute for Materials Research, Tohoku University, Sendai, Japan 2. The Institute of Scientific and Industrial Research, Osaka University, Ibaraki, Japan [email protected] Keywords: nanoparticle, chemical order, aberration correction
Hard magnetic properties of FePd nanoparticles originate from their L10-type tetragonal ordered structure with a high magnetocrystalline anisotropy energy. A correlation has been reported between the long-range order parameter and the particle size based on nanobeam electron diffraction [1]. However, the detailed nanostructure of small FePd nanoparticles is not fully understood. In this study, spherical aberration (Cs) corrected transmission electron microscopy (TEM) was employed for high-resolution observation of the atomic structure of FePd nanoparticles 2-10 nm in diameter. Besides its highly improved spatial resolution, a Cs-corrected TEM has a small optimal defocus condition due to a small Cs value (Δf = (4/3Csλ)0.5 ), which results in clear visualization of particle surface region. FePd nanoparticles were fabricated by successive deposition of Pd and Fe onto NaCl(001) substrates at 673 K [2]. After the deposition of Fe, an amorphous Al2O3 thin film was deposited to protect the particles from oxidation. Post-deposition annealing at 873 K led to a formation of the L10-type ordered structure. High-resolution TEM (HREM) images were obtained by using FEI Titan 80-300 TEM operated at 300 kV with a field emission gun and a Cs corrector for the objective lens. All TEM images were recorded by using CCD camera attached to the TEM. HREM images of FePd nanoparticles with c-axis oriented normal and parallel to the film plane are shown in Figures 1(a) and 1(b), respectively. Periodic arrangement of atoms due to chemical order is clearly seen as bright contrasts due to a small negative value of corrected Cs. Fourier transforms also indicated high spatial resolution. Particle periphery, i.e., the interface between crystal and amorphous, is also clearly observed. The analyses of power spectra of the amorphous region led us to conclude that the defocus value is about 13 nm, which is much smaller than the typical Scherzer defocus value (40-50 nm) for conventional high-resolution microscopes. A clear-cut long-range order is lost when particle size is smaller than about 5 nm, and partially ordered particles become dominant. It may be suggested that smaller optimal defocus is suitable for precise observation of atomic structure of nanoparticles. 1. 2. 3.
K.Sato, Y. Hirotsu, H. Mori, Z. Wang and T. Hirayama, J. Appl. Phys. 98 (2005), p.024308. K.Sato and Y. Hirotsu, J. Appl. Phys. 93 (2003), p.6291. This study was partially supported by the Grant-in-Aid for Young Scientists (B) (Grant No. 19760459) from the Ministry of Education, Culture, Sports, Science and Technology, Japan.
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Figure 1. HREM images of FePd nanoparticles with the L10 structure: (a) beam || [001], (b) beam || [100]. A schematic illustration of the L10 structure is shown in the inset.
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Structural and chemical characterization of Co-doped ZnO layers grown on Si and sapphire R. Schneider1, L.D. Yao1, D. Gerthsen1, G. Mayer2, M. Fonin2, and U. Rüdiger2 1. Laboratorium für Elektronenmikroskopie, Universität Karlsruhe, D-76128 Karlsruhe 2. Fachbereich Physik, Universität Konstanz, D-78457 Konstanz [email protected] Keywords: Co-doped ZnO, ferromagnetism, TEM, EFTEM
Ferromagnetism at room temperature has been experimentally observed for several doped oxidic semiconductors (Co:TiO2, Mn:TiO2, Co:ZnO, Mn:ZnO, etc.) [1], but its origin is still under discussion. Therefore, Co(3 at%)-doped ZnO layers of approximately 150 nm thickness were grown on (001)Si and (001)sapphire substrates at 500 °C by magnetron sputtering. For each substrate the Zn0.94Co0.06O layers were deposited under oxygen-poor and -rich conditions by using Ar and Ar/O2, respectively, for sputtering and post-annealing in O2 to ensure oxygen saturation for samples prepared under oxygen rich conditions. The magnetic properties of the Co-doped samples were investigated by SQUID (superconducting quantum interference device) magnetometry. Only samples that were grown under oxygen-poor conditions with no oxygen post-annealing on sapphire show a small ferromagnetic contribution to the paramagnetic background, while oxygen-rich ones and those on Si substrate do not, even at low temperatures. To correlate the magnetic behavior with both structural and chemical materials properties, transmission electron microscopy (TEM) including highresolution and analytical TEM (HRTEM/AEM) have been applied. In the case of silicon as substrate for both sample types, i.e. oxygen-poor and -rich ones, the Zn0.94Co0.06O layer showed a columnar growth and, moreover, a hexagonal wurtzite structure with predominantly c-axis orientation in growth direction. This is visualized by Fig. 1 representing an HRTEM image and a selected-area electron diffraction (SAED) pattern of the interface region between the Si substrate and Co:ZnO layer. The SAED pattern contains two sets of reflections related to ZnO, which can be indexed according to [110] and [210] zone axes, indicating two different azimuthal orientations. Most of the ZnO columnar grains grow close the [001] orientation, but the analysis of the arc of the (002) reflection yields an opening angle of about 11.5° arc (Fig. 1b) demonstrating the existence of slightly tilted grains. For the Zn0.94Co0.06O layers grown on sapphire (not shown here) two different orientation relationships between ZnO and the substrate were found as well. In both cases Co:ZnO [001] grows parallel to Al2O3 [001] with two azimuthal orientations: I) Co:ZnO [110] // sapphire [110] and Co:ZnO [210] // sapphire [210] with a mismatch of ~ 32 % between layer and substrate and II) Co:ZnO [210] // sapphire [110] and Co:ZnO [110] // sapphire [210] (~18 % mismatch), where the situation (II) can predominantly be found for Co:ZnO grains in neighborhood of Al2O3 and growth temperatures above ~500°C. For both types of substrates energy-filtered TEM (EFTEM) using the Co-L23 edge revealed a homogeneous Co distribution in the oxygen-poor Co:ZnO layer. In contrast, S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 621–622, DOI: 10.1007/978-3-540-85226-1_311, © Springer-Verlag Berlin Heidelberg 2008
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regions of about 50 nm in extension with Co enrichment were detected in the oxygenrich Zn0.94Co0.06O layer (Fig. 2d). Secondary Co phases like Co3O4 and CoO (except for the wurtzite structure) were not detected. However, CoO with the same crystal structure as ZnO cannot be excluded. These Co-rich grains are predominantly existent in the very interface region to the Si substrate, and only partly within the Co:ZnO layer. Since the oxygen-poor Co:ZnO layer on Si did not behave ferromagnetic, these regions enriched in Co cannot be the only reason for it. Obviously, there is an additional influence of the substrate, i.e. silicon or sapphire, respectively, on the magnetic behavior. Further investigations are in progress to reveal the correlation between structural and chemical peculiarities of the Co:ZnO layers under particular consideration of the substrate on one hand and magnetism on the other. 1. 2.
J. M. D. Coey, M. Venkatesan, C.B. Fitzgerald, Nature Mater. 4 (2005) 173. The funding of this work within the project C10 of the Competence Network “Functional Nanostructures” by the country Baden-Württemberg is gratefully acknowledged.
Co:ZnO
Si 5 nm
a)
b)
Figure 1. a) HRTEM image of the interface region between Si substrate and Zn0.94Co0.06O layer (oxygen-rich conditions), b) typical SAED pattern.
Co:ZnO
Co:ZnO a)
c)
Si
200 nm
b)
Si
200 nm
d)
Figure 2. TEM bright-field images (a, b) and corresponding energy-filtered images of the Co distribution (c, d) in Zn0.94Co0.06O layers on Si (left: oxygen-poor growth conditions, right: oxygen-rich).
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TEM studies of cobalt-doped zinc oxide films J. Simon1, K. Nielsen2, M. Opel2, S.T.B. Goennenwein2, R. Gross2 and W. Mader1 1. Institut für Anorganische Chemie, Rheinische Friedrich-Wilhelm Universität, Römerstraße 164, 53117 Bonn, Germany 2. Walther-Meissner-Institut, Bayerische Akademie der Wissenschaften, Walther-Meißner-Straße 8, 8574 Garching, Germany [email protected] Keywords: EFTEM, thin films, segregation
Diluted magnetic semiconductors (DMS) such as cobalt-doped zinc oxide have been studied intensively during the last years, because of their possible combination of semiconducting and magnetic properties. Cobalt-doped ZnO is a DMS material with reported TC well above room temperature [1,2]. However, in other studies no ferromagnetism could be observed [3]. Due to these contradictory experimental results the nature of the ferromagnetism as well as the mechanisms of exchange coupling is still under debate. In particular, the presence of ferromagnetic inclusions – such as metallic ferromagnetic clusters – in the ZnO host crystal could be a possible origin of ferromagnetism. Using a 300 kV field-emission transmission electron microscope (CM300 UT) equipped with an imaging electron energy filter (Gatan GIF) we have studied different Zn0.95Co0.05O films grown on (0001) ZnO substrates. Films were grown by using pulsed laser deposition and they shows ferromagnetic-like magnetization curves at 300K. Epitaxial cobalt-doped ZnO-films grown at 320°C show defects in the film typical for doped ZnO (Figure 1). Energy-dispersive X-ray measurements (EDX) result in a Co content of 6±1 at. % of the cations. In some regions of the film peculiar spherical structures are observed as shown in the HRTEM image (left) of Figure 2. Electron spectroscopic imaging (ESI) of these structures performed at the Co-L edge (right) reveals that the structural change is related to a significant enrichment of Co. Further ESI studies show that the Co enrichment goes along with a decrease of zinc and oxygen, respectively. Hence the structural irregularities are metallic Co-clusters within the Codoped ZnO films, which explain the magnetic properties. 1. 2. 3. 4.
M. Venkatesan, C. B. Fitzgerald, J. G. Lunney, and J. M. D. Coey, Phys. Rev. Lett. 93 (2004), p. 177206. K. Nielsen, S. Bauer, M. Lübbe, S. T. B. Goennenwein, M. Opel, J. Simon, W. Mader, and R.Gross, phys. stat. sol. (a) 203 (2006), p. 3581. S. Kolesnik, B. Dabrowski, and J. Mais, J. Appl. Phys. 95, (2004), p. 2582. We kindly acknowledge the Deutsche Forschungsgemeinschaft for financial support via SPP1157.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 623–624, DOI: 10.1007/978-3-540-85226-1_312, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. HRTEM image of a Co-doped ZnO film grown on ZnO substrate.
Figure 2. Left: HRTEM image with special structure marked by arrow. Right: corresponding Co – map shows enrichment of Co at the special structure.
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Nanocrystallization of amorphous Fe40Ni38B18Mo4 alloy D. Srivastava, A.P. Srivastava and G.K. Dey Materials science Division, Bhabha atomic Research Center, Mumbai 400094, India [email protected] Keywords: nanocrystals, softmagnetic material, TEM
Nanocrystallization of amorphous magnetic materials, in general, leads to improvement in soft magnetic properties. It shows best soft magnetic properties when amorphous matrix consists of nano-particles (<50 nm)[1]. A better understanding of evolution of the microstructure is required to control the magnetic properties, which includes the nature, kinetics, size and shape of the evolving crystalline phases and the defects. In this study Amorphous ribbons (~ 40 µm thickness) of Fe40Ni38B18Mo4 alloy composition were produced by melt spinning and Nanocrystalline phases were obtained by controlled heat treatment of amorphous precursor. The two crystallization peaks obtained in DSC at temperatures 414°C and 522°C corresponded to primary (PP) and secondary (SP) phases respectively. Kissinger approach, Ozawa method and Yi-Qun Gao method were employed to determine the kinetic parameters and results are presented in Table 1. XRD pattern showed melt spun ribbon to be amorphous and samples heat treated at 430oC as γ-(Fe, Ni) (PP) phase. Samples heat treated at 530oC and 600oC, consisted of cubic (Fe,Ni,Mo)23B6 (SP) phase in addition to the PP. In order to determine particle size and relative volume fraction of PP and SP, crystalline peak were fitted to Lorentzian function and amorphous halo was fitted using a Gaussian function. The volume fraction of SP was found to increase with increase in temperature but it decreased when treated for longer duration at 530oC as well as 600oC. Particle sizes of the PP and SP was found to be in the range of (3.5-13 nm) and (13-27nm) respectively. The coarsening rate of PP than SP was lower. Figure 1 shows representative TEM micrograph and SAD patterns obtained from samples heat treated at 430oC and 530oC for 72h. The characteristic diffuse halo pattern of ribbon confirmed it to be amorphous. In the case of 430oC sample, observation of crystalline ring pattern typical of ultra fine grains suggested formation of nanocrystals as seen in dark field micrograph. This phase could be indexed as γ-(Fe, Ni)(PP). The nanocrystals having size range of 3-6 nm were uniformly distributed in the amorphous matrix. Similar to XRD study, a careful examination of the sample in the TEM did not show presence of any secondary phase. The dark field micrographs of the samples treated at 530oC and 600oC for 72h (heat treated above secondary crystallization temperature) were showing bimodal distribution of the nanocrystals(size range 8-10 nm (smaller) and 15-20 nm (larger) 530oC for 72h). It consisted of γ-(Fe, Ni) (PP) and cubic (Fe, Ni, Mo)23B6 (SP) phases. The larger size crystals were seen to coarsen faster in sample treated at 600oC in comparison to the sample treated at 530oC. These morphological and structural observations are entirely consistent with the observations made in XRD study.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 625–626, DOI: 10.1007/978-3-540-85226-1_313, © Springer-Verlag Berlin Heidelberg 2008
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Ms and Hc values of samples heat treated at 430°C and 530°C for 10 min and 60 min were measured at room temperature using VSM. The hysteresis loops in all the crystallized samples were found to be narrow and rectangular in shape and result is presented in Table 2. Both Hc and Ms values were found to decrease after crystallization. The extent of decrease was found to be function of both time and temperature of treatment. Table 1
Table 2
φ (°C /m)
n
E (kJ/mol)
n
E (kJ/mol)
5 20
2.9 2.6
(252.5± 20.9)
1.9 1.8
(375.8± 55.7)
40
2.2
PP
SP
1.5
(a)
(b)
Sample
particle size
as received 430°C /10m 430°C /60m 530°C /10m 530°C /60m
11 12 18 19
Hc Ms (A/m) (emu/g) 50 17 15 19 17
105.45 85.36 75.01 88.48 75.91
(c)
0.5 μm
Figure 1 Electron micrographs and SAD patterns obtained from (a) melt spun ribbon (b) and(c) heat treated at 430oC and 530oC for 72h respectively 1.
G. Herzer, Hand Book of Magnetic Materials, Elsevier Science, Amsterdam, vol 10 (1997) chap. 3, p 415
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Structural and compositional properties of Sm-Fe-Ta magnetic nanospheres prepared by pulsed-laser deposition at 157 nm in a N2 atmosphere S. Šturm1, K. Žužek Rožman1, E. Sarantopoulou2 and S. Kobe1 1. Jožef Stefan Institute, Nanostructured materials, 1000 Ljubljana, Slovenia 2. Theoretical and Physical Chemistry Institute, National Hellenic Research Foundation, TCPI, Athens 11635, Greece [email protected] Keywords: nanomagnetism, AEM, nanospheres
Sm-Fe-Ta intermetallic alloys in their bulk form can exhibit hard magnetic properties after absorbing nitrogen at the interstitial sites within the 2:17 rhombohedral structure [1]. Lately, these properties have become attractive for the expanding field of nanotechnology, and much effort has been put into the synthesis of various nanostructures that will potentially serve as the building blocks in future electronic devices, such as circulators, magnetostrictive heads and nano-electro-mechanical systems (NEMS). One of the most promising routes for the synthesis of nanostructured Sm-Fe-Ta intermetallic alloys in the form of nanosized spheres appears to be pulsedlaser deposition (PLD), performed in a nitrogen background pressure using a target with the composition Sm13.8Fe82.2Ta4.0 [2]. There is, however, a lack of insight into the formation mechanisms, which hinders the prediction of the conditions under which nanostructures with better magnetic properties could be formed. For this reason we systematically studied the structure and the composition of Sm-Fe-Ta-based nanospheres synthesized by PLD. By applying high-resolution transmission electron microscopy (HRTEM - Jeol 2010F, Cs=0.5 mm), electron energy-loss spectroscopy (EELS - Gatan DigiPEELS 766) and energydispersive X-ray spectroscopy (EDXS- LINK ISIS EDS 300) we found that at least two distinct types of Sm-Fe-Ta-based spheres appear that differ in their crystallinity and composition. Two characteristic spheres are shown in a high-resolution TEM image (Figure 1a). The lower-left sphere, marked as (type I), is characterized by a crystalline central part of the sphere and outer 20-nm-thick amorphous rim. The crystalline core was confirmed by the Fast Fourier Transform (FFT) calculated from the central part of the sphere. Discrete diffraction points, which are observed in the FFT are due to the existence of crystal lattice fringes in the related section of the HRTEM image (Figure 1a). The FFT analysis performed from the central part of the upper-right sphere, marked as (type II), shows no apparent periodic structure, indicating that this sphere is most likely amorphous. Combined EELS and EDXS analyses of this sphere confirmed the presence of the target elements, i.e., iron, samarium, tantalum and a high concentration of oxygen, with the approximate composition of Sm23Fe25Ta2O50. Moreover, no significant evidence of nitrogen present in this sphere was obtained with the EELS analysis (Figure 1b). In contrast, nitrogen was detected in the central part of the spheres
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with the crystalline core (type I), as shown in Figure 1b, confirming that nitriding of the nanospheres had taken place already during the PLD process. In this case, the central part of the sphere contained a significantly lower amount of oxygen, whereas the outer amorphous rim of the sphere has a composition matching the composition of the amorphous sphere of the type II. Assuming that the oxygen measured in the central part originates from the oxidized amorphous shell surrounding the crystalline metallic core, the composition of the core would correspond to Sm9Fe80Ta1N10, which is reasonably close to the composition of their bulk counterparts. The measured coercivity (μ0HC) of the as-deposited sample was 20 mT, which implies the soft magnetic behaviour of the nanospheres. This behaviour could be explained by the fact that the soft magnetic rim represents the nucleation centres for the magnetization reversal of the whole sphere, and that the sphere without the soft rim would exhibit much higher coercivity. 1. 2.
J.M.D. Coey, H. Sun, J. Magn. Magn. Mater., 87, 1990, L251-L254 S. Kobe, E. Sarantopoulou, G. Dražić, J. Kovač, M. Janeva, Z. Kollia, A.C. Cefalas, Appl. Surf. Sci. 254 (4), 2007, p. 1027-1031 The authors acknowledge financial support from the European Union under the Framework 6 program under a contract for an Integrated Infrastructure Initiative. Reference 026019 ESTEEM.
3.
a)
Type I
b)
Fe-L2,3
Type II
Type I
Type I O-K N-K
Type II 20 nm
Type II 400
500
600 eV
700
800
Figure 1. a) High-resolution TEM image of Sm-Fe-Ta-based nanospheres marked as (Type I) and (Type II), with the superimposed FFT obtained from the central part of the corresponding spheres. b) The corresponding background-subtracted EEL spectra acquired from the central part of spheres.
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Characterization of Ni-Mn-Ga magnetic shape memory alloys using electron holography and Lorentz microscopy K. Vogel, M. Linck, Ch. Matzeck, A. Rother, D. Wolf and H. Lichte Triebenberg-Labor, Institut für Strukturphysik, Technische Universität Dresden, 01062 Dresden, Germany [email protected] Keywords: electron holography, Lorentz microscopy, magnetic shape memory alloys
Ni-Mn-Ga magnetic shape memory alloys are considered to have a high potential for new actuating devices and sensors. The large magnetic-field-induced strain (6-9%) observed in these alloys is caused by magnetic-field-controlled twin boundary motion [1], [2]. For understanding the mechanism of twin boundary motion and the influence of microstructural parameters like grain size, texture, defects, structure of the martensite, a comprehensive characterization of microstructure and magnetic structure is required. Here we report on application of Electron Holography and Lorentz Transmission Electron Microscopy for characterization of polycrystalline Ni-Mn-Ga samples. The conventional in-focus bright-field image shows the twin-band microstructure of the martensitic material (Figure 1a), while the strongly defocused Lorentz image of the same sample area (Figure 1b) additionally shows the magnetic domain walls as bright and dark lines. There is evidently a strong correlation between the twin-bands and the magnetic domains. The twin-boundaries coincide with magnetic 90°-domain walls. However, the twin-bands are not single magnetic domains, instead they are subdivided into several magnetic domains by 180°-domain walls. In order to get more detailed information on the magnetization distribution inside the domains, we perform electron holography. The equi-phase lines in the phase image (Figure 2b) display the 3D magnetization distribution projected into the recording plane. The comparison of the experimental phase image, which was reconstructed from the hologram acquired within the broad twin-band marked in Figure 2a, with a simulated phase image for a wedge shaped sample containing 180° domains, gives a good qualitative agreement (Figure 3). Further work is in progress to eliminate the phase shift contribution stemming from thickness variation, to enable a quantitative analysis of the magnetization distribution.
1. 2. 3.
S. J. Murray, M. Marioni, S. M. Allen, R. C. O´Handley, T. A. Lograsso, Applied Physics Letters 77 (2000), p. 886. A. Sozinov, A. A. Likhachev, N. Lanska, K. Ullakko, Applied Physics Letters 80 (2002), p. 1746. We thank Dr. Andrea Böhm, Fraunhofer IWU Dresden, and Dr. Stefan Roth, IFW Dresden, for the specimens. Financial support from Deutsche Forschungsgemeinschaft within DFG Priority Programme 1239 is gratefully acknowledged.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 629–630, DOI: 10.1007/978-3-540-85226-1_315, © Springer-Verlag Berlin Heidelberg 2008
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a
b
c
Figure 1. a) In-focus bright-field image of twin-bands, twin-boundaries marked by arrows. b) Lorentz image (underfocus) showing magnetic domain walls as bright and dark lines. c) Subimage from image b) revealing interference fringes in the bright lines.
a
b
va cu um
vacuum
TB
TB
Figure 2. a) Lorentz image (overfocus) showing the position of hologram acquisition marked by the box; twin boundaries (TB) are marked by arrows. b) Reconstructed phase image; equi-phase lines mod(π/2) give a coarse impression about the distribution of the projected in-plane component of the magnetic B-field lines, as indicated by the arrows. wedge shaped sample with 180° domains [nm]
simulated phase image mod(2π)
experimental phase image mod(2π)
vacuum
[nm]
Figure 3. Simulated phase image for a wedge shaped specimen containing 180° domains, and comparison with the experimental phase image obtained from the broad twin-band shown in Figure 2a.
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Energy Loss Magnetic Chiral Dichroïsm (EMCD) for magnetic material B. Warot-Fonrose, L. Calmels, C. Gatel, F. Houdellier, V. Serin and E. Snoeck CEMES-CNRS, 29 rue Jeanne Marvig, 31055 TOULOUSE, France [email protected] Keywords: EMCD, magnetic materials, EELS
Trends in electronics propose the use of spin instead of charge as an information unit. These new spintronic devices are based on magnetic and non-magnetic thin layers through which electrons travel and are scattered or filtered at interfaces depending on their spin. The spin mean free path depends on the material but can be estimated to be in the 100 nm range. The interesting devices for spintronics have therefore dimensions close to few hundred of nanometer and therefore physical measurements have to be performed at this scale. One of the important properties is the local magnetic moment and its evolution close to surfaces or interfaces. X-ray magnetic circular dichroism has been a powerful technique to measure magnetic moments on µm size areas. However new set-ups need to be built now to take into account the reduced size of the regions of interest. Experimental techniques capable of measuring the magnetic moment at this scale are developed, especially photo emission electron microscopy (PEEM, [1]). This technique is very promising but needs synchrotron radiation. The nanometer size of the probe in a transmission electron microscope (TEM) is compatible with the resolution needed for magnetic measurements on spintronic devices. Hébert et al. [2] proposed an original set-up to combine the precise location of the probe in a TEM and the measurement of magnetic properties. This technique, called energy-loss magnetic chiral dichroism EMCD [3], is based on the distribution in the diffraction plane of the inelastically scattered electrons that have crossed the magnetic sample. The signal is detected with an electron energy loss spectrometer to probe the inelastic interactions. The major issue in these experiments is the very low signal. Therefore experimental acquisitions have to be optimized to get a reliable signal. We have proposed any original set-up using the spherical aberration corrector and the energy spectrum imaging technique [4], we have also developed a quantitative analysis derived from the X-ray sum rules [5]. We will detail this technique and its application to iron and magnetite thin films. These materials have been chosen to improve our experimental set-up and develop the correction methods needed to extract reliable data from the spectra [6]. This method is used to the measure the modification of the iron magnetic moment when the atoms have an environment different from the bulk iron. We have studied FexCo1-x alloys because magnetic measurements and ab-initio calculations have shown that the magnetic moment of iron depends on its chemical environment. Results on Fe/Co multilayers will also be presented to show the role of the interfaces.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 631–632, DOI: 10.1007/978-3-540-85226-1_316, © Springer-Verlag Berlin Heidelberg 2008
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1. 2. 3. 4. 5. 6.
J. Stöhr, Y. Wu, B.D. Hermsmeier, M.G. Samant, G.R. Harp, S. Koranda, D. Dunham, and B.P. Tonner, Science 259 (1993), p.658 C. Hébert, P. Schattschneider, S. Rubino, P. Novak, J. Rusz and M. Stöger-Pollach, , Ultramicroscopy (2007), doi:10.1016/j.ultramic.2007.07.011 P. Schattschneider, S. Rubino, C. Hébert, J. Rusz, J. Kunes, P. Novak, E. Carlino, M. Fabrizioli, G. Panaccione, G. Rossi, Nature 441 (2006), p.486 B. Warot-Fonrose, et al., Ultramicroscopy (2007), doi:10.1016/j.ultramic.2007.05.013 L. Calmels, F. Houdellier, B. Warot-Fonrose, C. Gatel, M. J. Hÿtch, V. Serin, E. Snoeck, and P. Schattschneider, Phys. Rev. B, 76 (2007), p.060409 B.Warot-Fonrose, C.Gatel, F.Houdellier, P.Schattschneider, Mat. Res. Soc. Symp. Proc., 1026E, (2007), C13-06
a)
b)
1
2
pos 1 pos 2
c)
690
700
710
720
730
740
energy (eV)
Figure 1. a) ESI acquisition of the intensity of the diffraction pattern on a Fe single film, b) diffraction pattern and drawing of the Thales circle with the 1 and 2 positions used to measure the dichroic signal, c) Spectra acquired in position 1 and 2 on the corrected data cube
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Determining the nanoscale chemistry and behavior of interfaces and phases in Al-Si(-Cu-Mg) nanoparticles using in-situ TEM J.M. Howe1, S.K. Eswaramoorthy1 and G. Muralidharan2 1. Department of Materials Science and Engineering, University of Virginia Charlottesville, VA, 22904-4745 USA 2. Materials Science and Technology Division, Oak Ridge National Laboratory Oak Ridge, TN 37831-6132 [email protected] Keywords: solid-liquid interface, nucleation, in-situ TEM
Understanding fundamental phenomena such as the structure and composition of the solid-liquid interface, the partitioning of elements between the solid and liquid phases during crystal growth, and critical factors involved in the nucleation of phases, are essential to solidification science and the development of novel material structures [1]. The sample chosen for this investigation was a commercial Al-Si base alloy (AA390) powder containing Al-17.87Si-1.83Cu-0.6Mg (at.%). The powder particles were suspended in ethanol and dispersed on an ultra-thin carbon film supported on a coppermesh TEM grid. In-situ heating experiments were conducted at 200 kV in a JEOL 2000FX-II TEM using a Gatan double-tilt heating holder. Energy-dispersive X-ray spectroscopy (EDXS) was performed using a Gresham high-angle X-ray detector and an incident electron beam diameter of about 25 nm [2,3]. Figure 1a shows a bright-field TEM image of an Al-Si-Cu-Mg powder particle about 350 nm in diameter, taken at 828 K. The particle is partially molten, containing a faceted primary Si particle on the lower left, and an oval-shaped solid Al-rich particle surrounded by Al-rich liquid phase in the upper-right. X-ray spectra were obtained at five positions indicated by circles a-e in Figure 1a and the resulting compositions are plotted as solid lines in Figures 1c and d. From these figures, it is evident that the composition on the right side of the particle, position e, contains approximately the same amounts of Cu and Mg as that in the Si on the far left, position a. In addition, as the thin, liquid interface between the two solid phases is approached from either side, the Al and Si concentrations change in a complementary and symmetric manner. Figure 1b shows the same particle after further heating to 878 K. The particle now contains only two phases, solid Si in the lower-left and Al-rich liquid. X-ray spectra were obtained at four positions indicated by circles 1-4, and the resulting compositions are plotted as dashed lines in Figures 1c and d. As at 828 K, the Al and Si concentrations change in a complementary and symmetric manner about the solid-liquid interface, although the Al concentration in the liquid on the far right has decreased while the levels of Cu and Si have increased. In contrast to 828 K, the Cu and Mg concentrations both decrease in going from the liquid into the solid Si at 878 K.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 633–634, DOI: 10.1007/978-3-540-85226-1_317, © Springer-Verlag Berlin Heidelberg 2008
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In addition to EDXS, conventional imaging enables direct observation of the behavior of the various phases. The presence of the liquid phase surrounding the solid Al phase in Figure 1a indicates that it is energetically unfavorable for the Al-rich solid to form an interface with either the solid Si or the oxide shell (i.e., the contact angle is 180º). This provides direct evidence that neither of these phases could have acted as a catalyst in the nucleation of the Al-rich solid, in agreement with Das et al. [4].
Figure 1. (a) Bright-field TEM image of a partially molten Al-Si-Cu-Mg powder particle taken at 828 K. The faint, dashed line parallel to the interface in (a) indicates the approximate appearance of the column of material sampled by the electron beam normal to the plane of the figure. (b) Bright-field TEM image of the same powder particle after further heating to 878 K. (c) Concentration profiles of Al and Si obtained from X-ray spectra taken at 828 K (solid lines) and at 878 K (dashed lines). (d) Concentration profiles of Cu and Mg obtained from X-ray spectra taken at 828 K. 1. 2. 3. 4. 5.
M. Rappaz, C. Beckermann and R. Trivedi, eds., "Solidification Processes and Microstructures: A Symposium in Honor of Prof. W. Kurz", (TMS, Warrendale) (2004). S. K. Eswaramoorthy, J. M. Howe and F. Phillipp, Micros. Microanal. 13 (2007) p. 291. S. K. Eswaramoorthy, J. M. Howe and G. Muraldiharan, Science 318 (2007) p. 1437. S. K. Das, J. H. Perepezko, R. I. Wu and G. Wilde, Mater. Sci. Eng. A304 (2001) p. 159. This research was supported by NSF under Grant DMR-0554792 and the U.S. Department of Energy under contract DE-AC05-00OR22725 with UT-Battelle, LLC.
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Dislocations in AlPdMn quasicrystals: contrast in TEM and physical properties D. Caillard, F. Mompiou CEMES, CNRS, 29 rue Jeanne Marvig, BP4347, 31055 Toulouse Cedex, France [email protected] Keywords: quasicrystals, AlPdMn, dislocations, phason walls, bright-field contrast
Defects in quasicrystals are simpler than what could be anticipated from the rather exotic quasicrystalline structure. For instance, dislocations exhibit properties (contrast in TEM, possible modes of motion under stress) that are close to those in crystals. The aim of this talk is to introduce these properties using simple concepts, and avoiding complex calculations in high-dimensional space. The structure of quasicrystals will be first described starting from a 2-dimensional aperiodic tiling obtained by the cut and projection of a 3-dimensional periodic cubic lattice along an irrational plane. Perfect (retiled) and imperfect (non-retiled) dislocations will be defined in this 2-dimensional quasicrystal. Then, the icosahedral structure of AlPdMn will be defined by the same procedure, namely by the cut and projection of a 6dimensional long-range-ordered hypercubic lattice [1]. Since coherent diffraction can take place in spite of the lack of periodicity, dislocations and stacking faults exhibit TEM contrasts that can be compared to those in crystals. Different examples will be given of the bright-field contrast of perfect dislocations and imperfect dislocations trailing phason planes (Figure 1). For perfect dislocations, this contrast can be interpreted only when the phase shifts due to phason fields (local chemical disorder) and elastic strain fields are both taken into account. Since the phason field can be described by a displacement in the “perpendicular space” lost during the cut and projection process, rules of contrast are the same as in normal crystals provided the scalar products G.B are computed in the 6-dimensional periodic space. Superdislocations dissociated into two superpartials are separated by a complex fault which can be described as the superposition of an antiphase boundary and a phason wall. Recent results showing that climb is the dominant mechanism of dislocation motion will be presented [2, 3]. In particular, a high-temperature in situ experiment shows that glide is at least 1000 times slower than climb [4]. This unusual property has been directly related to the quasicrystalline structure. 1. 2. 3. 4.
F. Mompiou and D. Caillard, Materials Science and Engineering A400-401 (2005) p. 283. F. Mompiou, L. Bresson, P. Cordier, and D. Caillard, Philosophical Magazine 83 (2003) p. 3133. F. Mompiou, D. Caillard and M. Feuerbacher, Philosophical Magazine 84 (2004) p. 2777. F. Mompiou and D. Caillard, Philosophical Magazine Letters 84 (2004) p. 555.
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Figure 1. Bright-field images of dislocations and phason walls, in AlPdMn deformed at 300°C. From [1]
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Characterization of a-plane GaN films grown on r-plane sapphire substrate by electron microscopy Y. Arroyo Rojas Dasilva1, T. Zhu2, D. Martin2, N. Grandjean2 and P. Stadelmann1 1. Interdisciplinary centre for electron microscopy, Ecole Polytechnique Fédérale de Lausanne, CH-1015 Lausanne, Suisse. 2. Laboratory of advanced semiconductors for photonics and electronics, Ecole Polytechnique Fédérale de Lausanne, CH-1015 Lausanne, Suisse. [email protected] Keywords: a-GaN, dislocations, stacking faults
Gallium nitride (GaN) is a semiconductor material which has recently attracted much attention for wide and direct band gap of 3.39 eV. It has been studied intensively over the last years due to the possible applications in optoelectronic and electronic devices such as light emitting diodes [1], laser diodes [2], which are active in the green, blue and ultraviolet wavelength, detectors [3] and high temperature/high power electronics [4]. C-plane GaN is the most common orientation to grow GaN devices. However, GaN suffers of polarization effects originating from c-[0001] polar axis of the wurtzite structure. These effects produce strong electric fields that affect the lifetime and efficiency of the devices. One approach to avoid this problem is to grow GaN as non polar structures (m-{1-100}[5] and a-{11-20} [6]) planes GaN. There is a considerable interest to analyze the defects in a-plane GaN layers and their effect in the optoelectronic properties of GaN-based devices that are grown on r-plane sapphire in non polar directions. This work presents the structural characterization by electron microscopy of a-plane GaN grown on r-plane sapphire. A-plane GaN films on r- plane sapphire were grown by hydride vapour phase epitaxy (HVPE)–epitaxy lateral overgrowth (ELO). ELO allows reducing the defects density and improving the electrical and optical properties. The GaN layer grows vertically in the [11-20] direction and laterally in the [0001] and [000-1] directions. Transparent cross sectional samples of GaN films for TEM were prepared by dimpler method and ion milling with Ar+ at 2.5 kV. The characterization of the a-plane GaN film consists in the identification of structural defects, mainly dislocations and stacking faults. The main defects in a-plane GaN layers are threading dislocations (TDs) and stacking faults (SFs). High defects density in a-GaN layers (Figure 1a and b) was found, the dislocation density is in the range of 1010-1011 cm-2 and stacking fault density is 104-105 cm-1. Threading dislocations in a-GaN layers are originated at the a-GaN/r-sapphire interface for the difference between the lattice parameters. They have an edge, screw and mixed character with a Burgers vector b=1/3<11-20> and b=1/3<1-213>, respectively. Partial dislocations are also found in the a-GaN layers, they are formed by the splitting of perfect dislocations that are associated with stacking faults. They are Frank dislocations with b=1/6<20-23> and Shockley dislocations with b= 1/3<1-100>. Intrinsic stacking
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faults of type I1 and I2 on the basal plane (Figure 1b) were found in the a-plane GaN layers with a displacement vector R=1/6<20-23> and R=1/3<1-100>, respectively . Some of them are originated at the interface GaN/sapphire or GaN/SiO2 interface and go through the layer. The others are overlapping the staking faults that are originated at the interface. A different stacking fault (PSFs) was also found in a-GaN (Figure 1b), prismatic stacking fault with R=1/2<10-11> on the (11-20) prismatic plane. 1. 2. 3. 4. 5. 6.
Chitnis, C. Cheng, A. Adivarahan, M, Shatalov, E. Kuokstis and M. Asif Khan. Appl. Phys. Lett 84 (2004), p.3663. S. Nakamura. J. Crystal Growth 201-202 (1999), p.290. E. Monroy, F. calle, J.L. Pau, E. Munoz, F. Omnes, B. Beaumont and P. Gibart. Phys. Stat. Sol. A 185 (2001),p.91. S. J. Pearton, F. Ren, A. P. Zhang and K. P. Lee. Mater. Sci. and Eng. R30 (2000) p.55. S. Ghosh, P. Waltereit, O. Brandt, H. T Grahn and K. H. Ploog. Phys. Review B 65 (2002), p. 0752021. D.S. Lee, H. Chen, H.B. Yu, X. H. Zheng, Q. Huang and J. M. Zhou. Chin. Phys. Lett, 21 (2004), p. 970. [11-20]
[0001]
[1-100]
ELO area
TDs and SFs opening
a)
b)
Figure 1. Structural defects in the a-GaN films: a) distribution of defects in the film, b) basal stacking faults (BSF) and prismatic stacking faults (PSF) in ELO area.
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A method for atomistic/continuum analysis of defects in large HRTEM images A. Belkadi1, G.P. Dimitrakopulos2, J. Kioseoglou2, G. Jurczak1, T.D. Young1, P. Dluzewski1, and Ph. Komninou2 1. Department of Computational Science, Institute of Fundamental Technological Research, Polish Academy of Sciences, Swietokryska 21, 00-049 Warsaw, Poland 2. Solid State Physics Section, Department of Physics, Aristotle University of Thessaloniki, GR-54124 Thessaloniki, Greece [email protected] Keywords: HRTEM, atomic structure, dislocation.
The capability to extract continuous lattice distortion fields from High Resolution transmission Electron Microscopy (HRTEM) images is invaluable for treating a large number of material problems, and methods such as Geometrical Phase Analysis (GPA) or peak-finding have been employed for this purpose. However, it is difficult to use such methods to obtain experimental strain maps from HRTEM images of large areas containing many dislocations due to localized changes of the imaging conditions imposed by the strain fields or by variations of the specimen thickness “Figure 1.a”. In this work we propose another method for the investigation of dislocation structures contained in large areas of HRTEM images and their strain maps. In the first step, the atomic structure including all dislocations visible on the HRTEM image is generated “Figure 1.b”. Independently of the atomic structure the maps of the lattice distortion tensor are then generated and the stress equilibrium configuration is determined by using a nonlinear finite element method for anisotropic nonlinear elastic structure. The stress/force equilibrium configuration of the lattice is determined both for the respective continuum problem and for the atomic structure. Using the atomic structure in a state of self-equilibrium a HRTEM simulated image is generated “Figure 1.c” and compared with the experimental one. By the method proposed here, problems induced on the HRTEM observation due to lattice curvature and foil effects can be accounted for and analyzed. Also the residual stress induced by these dislocations embedded in an anisotropic non-linear elastic structure is calculated. Compared to two-dimensional approaches such as the GPA and peak-finding, this method has the advantage of yielding three-dimensional fields surrounding the dislocations, which is significant in the case of mixed-type dislocations. Furthermore, this method conveniently overcomes the `size-problem' in the analysis of HRTEM images as it is applicable to large volumes of material with many dislocations trapped inside. 1.
Financial support for this work from the Marie Curie Research Training Network PARSEM (MRTN-CT-2004-005583) is gratefully acknowledged.
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(a)
(b)
(c) Figure 1. (a) HRTEM image of Gallium Nitride along [0001] showing four mixed-type threading dislocations in a large (29.3nm × 29.3nm) area of sample. (b) Threedimensional atomic structure with dislocations generated under the present methodology on the basis of the HRTEM image. (c) Simulated HRTEM image obtained from the 3D atomic structure.
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High resolution electron microscopy of interfaces in ultrafine microstructures of Zr and Ti based alloys G.K. Dey, S. Neogy, R.T. Savalia, R. Tewari, D. Srivastava and S. Banerjee Materials Science Division, Bhabha Atomic Research Centre, Trombay, Mumbai 400085, India [email protected] Keywords: metallic glass, nanocrystals, HREM
Nanosized grains have been produced in Zr based bulk metallic glasses and rapidly solidified metallic glasses by crystallization. The various types of interfaces generated in these microstructures have been examined by high resolution transmission electron microscopy (HREM). The nanocrystals produced by the crystallization of Zr54.5Cu20Al10Ni8Ti7.5 and Zr52Ti6Al10Cu18Ni14 bulk glasses have been found to lie in the size range of 15 to 50 nm and comprised phases isostructural with tetragonal Zr2Cu and Zr2Ni. These observations are similar to those made by Xing et al [1]. HREM examination of nanograin boundaries showed the extension of the lattice fringes right up to the grain boundary. The lattice could be resolved simultaneously in some of the grains in view indicating that grain boundaries are parallel to low index planes in these grains [2]. Crystallographic defects like stacking fault and antiphase domain boundary could be noticed in many of the nanograins. Figure 1 shows the high resolution image of a nanograin of Zr2Ni phase where besides the fundamental lattice fringes a domain like structure could also be revealed. A number of planar crystallographic faults could be observed in this nanograin. A detailed analysis of these crystallographic faults has been carried out in this study. The same alloys have been produced in the glassy state by rapid solidification. Crystallization of the rapidly solidified metallic glasses also led to the formation of nanocrystals where HREM examination revealed the presence of twins and twin-twin interaction. It has been possible to generate a very fine lamellar microstructure by combustion synthesis and by direct laser fabrication technique of TiAl with widths of the α2 and γ lamellae in the range of 5 to 20 nm. The interface between the α2 and γ phases has been investigated in detail in this microstructure (Figure 2). The interfaces between contiguous lamellae of the γ phase have also been examined. The interface across which one phase is transforming to the other could be deciphered by HREM. 1. 2.
L.Q. Xing, J. Eckert, W. Loser, L. Schultz and D.M. Herlach, Philosophical Magazine A 79 (1999), p. 1095. D.B. Williams and C.B. Carter in “Transmission electron microscopy III”, (New York; Plenum Press) 1996, p. 459.
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3.5 nm
Figure 1. HREM image of a nanocrystal of Zr2Ni phase in crystallized Zr54.5Cu20Al10Ni8Ti7.5 bulk metallic glass where many planar crystallographic faults can be seen within the nanograin.
3.3 nm
Figure 2. HREM micrograph showing the lamellar microstructure comprising of α2 and γ phases, the α2/γ and the γ/γ interfaces, faults in the α2 and γ phases.
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Anisotropic strain relaxation in (110) La2/3Ca1/3MnO3 thin films S. Estrade1, I.C. Infante2, F. Sanchez2, J. Fontcuberta2, J. Arbiol1,3, F. Peiró1 1. EME/CeRMAE/IN2UB, Dept. d'Electrònica, Universitat de Barcelona, c/ Marti Franques 1, 08028 Barcelona, CAT, Spain 2. Institut de Ciència de Materials de Barcelona -CSIC, 08193 Bellaterra, CAT, Spain 3. TEM-MAT, Serveis Cientificotecnics, Universitat de Barcelona, 08028, Barcelona, CAT, Spain [email protected] Keywords: LCMO, STO, TEM
Ferromagnetic metallic manganites, such as La2/3Ca1/3MnO3 (LCMO), are employed as electrodes in magnetic tunnel junctions (MTJs), which are heterostructures using ferromagnetic electrodes and insulating nanometric barriers. Different effects arising at the bare electrode or at the interfaces can affect the physical properties of the final device. Among them we highlight here the elastic deformation due to lattice parameter mismatch, the plastic strain relaxation, and possible stoichiometry changes through interdiffusion or segregation of the species. It has been observed that when using manganites in MTJs, the measurend magnetoresistane drops to zero at a temperature well below the Curie temperature of the bulk material. It has been recently shown that the magnetic properties of, simultaneously grown, epitaxial manganite films of different texture largely differ [1]. It has been suggested that this intriguing observation is related to anisotropic elastic properties of the manganite latice, although microstrutural analysis is not yet reported. In the present work we present transmission electron microscopy (TEM) and Electron energy loss spectroscopy (EELS) characterization of 37 nm LCMO layers grown on (110) STO substrates compared to those grown on (001) STO substrates. XRD measurements showed that [110] LCMO is further relaxed than [001] LCMO for any given thickness (in the studied 10-150 nm range) and that an anisotropic relaxation occurs, with [1-10] direction relaxing faster than ([001] [1]. TEM cross section (XT) observation of the layers (Figure 1) revealed [110] LCMO presents an abrupt interface but a somewhat rougher free surface when compared to its [001] counterpart, as well as the presence of some disoriented grains and of plastic defects in the layer. EELS experiments showed the layers to be chemically homogeneous. Plane view (PV) TEM characterisation gave further insight into the nature of the encountered defects. g = (1-10) two beam observation (Figure 2a) showed a periodic strain contrast at the interface position, corresponding to a highly periodic network of misfit dislocations running perpendicular to the [1-10] direction, with a spatial periodicity of about 50-100 nm. From XRD lattice parameter measurements, about 0.6% strain was found in the [1-10] direction and, in order to accommodate the resulting misfit, misfit dislocations with an average separation of 70 nm were expected along the [001] one. On
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the other hand, g = (001) two beam observation showed a much lower density of misfit dislocations, separated by some 50-300 nm, running perpendicular to (001) and extending to lengths below 100 nm (Figure 2b). XRD lattice parameter measurements yielded about 1.0 % strain in the [001] direction and, thus, mean separation between misfit dislocations was expected to be about 250 nm in this case. For both in-plane directions the distance between dislocations as measured from two beam experiments is in good agreement with the expected degree of relaxation. 1.
I.C. Infante, F. Sánchez, J. Fontcuberta, M. Wojcik, E. Jedryka, S. Estradé, F. Peiró, J. Arbiol, V. Laukhin, J. P. Espinós, Phys. Rev. B 76 (2007) 224415-1
Pt Au LCMO
STO
Figure 1. g=(1-10) two beam image of [110] LCMO/STO in XT geometry. Au and Pt correspond to protective layers in FIB sample preparation.
(b)
(a) LCMO + STO
LCMO + STO
g = (001) g = (1-10)
0.2 µm
Figure 2. g=(1-10) (a) and g = (011) (b) two beam images of [110] LCMO/STO in PV geometry.
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Metadislocations in complex metallic alloys: core structures investigated by aberration-corrected scanning transmission electron microscopy M. Feuerbacher, L. Houben, and M. Heggen IFF-IMF, Forschungszentrum Juelich GmbH, 52425 Juelich, Germany [email protected] Keywords: complex metallic alloys, metadislocations, dislocation core structure
Complex metallic alloys (CMAs) are materials of high structural complexity exemplified by a large number of atoms per unit cell and icosahedral-symmetric atomic coordination [1]. As a direct result of their large lattice parameters, conventional dislocation-based plastic-deformation mechanisms are prone to failure in these materials. Indeed, in several CMA phases, deformation mechanisms involving a novel defect type, the metadislocations, were identified [2]. In this contribution an investigation of metadislocation core structures by means of aberration-corrected high-resolution scanning transmission electron microscopy is presented. We have employed a FEI Titan 80 − 300 microscope equipped with a double-hexapole probe corrector and a high-angle annular dark-field (HAADF) detector. The latter allows for a direct and unambiguous characterization of structural defects in terms of a tiling representation, as well as for direct conclusion on the atom coordinations at the dislocation core. The CMA phases considered are ξ’-Al-Pd-Mn and T-Al-Mn-Pd, orthorhombic phases with 318 and 156 atoms per unit cell, respectively. In ξ’-Al-Pd-Mn we investigated metadislocations with different Burgers-vector magnitude. For the first time, we have resolved the atom coordination at the dislocation core and found that the strain field can be described in terms of overlapping substructure clusters in arrangements not found in the ideal structure (Fig. 1). The dislocations in T-Al-Mn-Pd could also be identified to be of metadislocation type. In contrast to the metadislocations in ξ’-Al-Pd-Mn, however, the core in the T-phase is accommodated into the structure by characteristic ancillary defects. The latter are shown to be integral parts of the dislocation and therefore an extraordinarily high number of atoms is involved in dislocation motion. 1. 2.
K. Urban and M. Feuerbacher, J. Non Cryst. Sol. 334 – 335 (2004), p.143. H. Klein, M. Feuerbacher, P. Schall and K. Urban, Phys. Rev. Lett. 82 (1999), p. 3468.
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1 nm Figure 1. HAADF –STEM micrograph of a metadislocation in ξ’-Al-Pd-Mn with a superposed tiling representation and cluster distribution.
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TEM of high pressure torsion processed intermetallic Zr3Al D. Geist, C. Rentenberger and H.P. Karnthaler Physics of Nanostructured Materials, Faculty of Physics, University of Vienna Boltzmanngasse 5, 1090 Vienna, Austria [email protected] Keywords: Zr3Al, severe plastic deformation, crystallographic boundaries
The ordered L12-structured intermetallic compound Zr3Al is reported to show a strong tendency to amorphization under ion- and electron-irradiation, hydrogenation and severe plastic deformation (SPD) by ball-milling [1]. In this work, Zr3Al has been investigated after SPD by high pressure torsion (HPT). Transmission electron microscopy (TEM) methods were used to elucidate the structural evolution of the compound at different stages of deformation by HPT. With a TEM operating at 200kV, images and diffraction patterns were taken at shear strains up to 140 000%. Since Zr3Al cannot be produced as a single-phase compound, energy dispersive X-ray measurements were performed to ensure that the composition in the investigated regions is Zr-25at.%Al. Figure 1a shows Zr3Al deformed to a shear strain of ~14 000%. Grain fragmentation occurs along {111} planes by twinning and by accumulation of superlattice intrinsic stacking faults (SISF). In the diffraction pattern (Figure 1b), in addition to the matrix spots (index m), diffraction spots of the twinned structure (index t) are visible. The (111)-plane is the twinning plane. SISF on the {111} planes in matrix and twin generate streaks in the diffraction pattern. Despite the high grade of deformation, superlattice reflections are still visible which means that the order is preserved. Figure 2 shows Zr3Al deformed even much higher (~140 000%). Grain fragmentation occurs along {111} planes; the orientation variation of the fragments is small as revealed by the diffraction pattern. Streaks in both <111> directions in the diffraction pattern imply crystallographic boundaries formed by accumulated SISF. Superlattice reflections are marked; they show that order is still present. Figure 2b shows a band confined by two different {111} planes. Due to the superposition of the orientations of the {111} planes, the band itself is oriented along a high index direction. The band shows complex contrast caused by a high defect density and Moiré patterns. Despite the high grade of deformation, no tendency to disordering, nanocrystallization or amorphization has been encountered. This contrasts very much to the results of HPT-processed L12-structured Ni3Al [2] and ball-milled Zr3Al [3], where a disordered nanocrystalline and an amorphous structure occured, respectively. 1. 2. 3.
E.M. Schulson in “Intermetallic Compounds - Structural Applications of Intermetallic Compounds”, ed. J.H. Westbrook, (John Wiley & Sons, Sussex) (1995), p. 137 C. Rentenberger, H.P. Karnthaler, Int. J. Mat. Res 98 (2007), p. 255 T. Benameur, A.R. Yavari, J. Mater. Res.7 (1992), p. 2971
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4.
Acknowledgments: The authors thank Prof. E. M. Schulson for the provision of Zr3Al, the group of Prof. R. Pippan for their kind help with the HPT deformation and for the support by the research project “Bulk Nanostructured Materials” within the research focus “Materials Science” of the University of Vienna. D.G. acknowledges the support by the I.K. “Experimental Materials Science – Nanostructured Materials”, a college for PhD students at the University of Vienna.
a)
b)
Figure 1. Zr3Al deformed by a shear strain of ~14 000% (BD ~ [01-1]). (a) Bright field image; the crystallographic boundaries at both sides of the (111)-twin boundary (TB) consist of accumulated SISF on {111}-planes. (b) Diffraction pattern; several superlattice reflections are marked by arrows.
a)
b)
Figure 2. Dark field images of Zr3Al deformed by a shear strain of ~140 000% (BD ~[01-1]) . (a) Grain fragmentation occuring on {111} planes. (b) Deformation band lying along a high index direction and bounded by steps parallel to {111} planes.
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Multiscale characterisation of the plasticity of Fe-Mn-C TWIP steels H. Idrissi1, L. Ryelandt2, K. Renard2, S. Ryelandt2, F. Delannay2, D. Schryvers1, P.J. Jacques2 1. EMAT, University of Antwerp, Department of Physics, Groenenborgerlaan 171, 2020 Antwerp, Belgium 2 .Université catholique de Louvain, Département des Sciences des Matériaux et des Procédés, IMAP, Place Sainte Barbe 2, B-1348 Louvain-la-Neuve, Belgium [email protected] Keywords: TWIP Steels, stress-strain response, twinning mechanism, dislocations, EBSD, TEM.
It is known for a long time that mechanical twinning occurring in FCC metal can bring about a very large work hardening rate. It is believed that deformation twins increase the work-hardening rate by acting as obstacles for gliding dislocations. This mode of plastic deformation is active, among others, in the Hadfield steel known for a very long time [1]. Fe-Mn-C grades exhibiting mechanical twinning present a renewed interest for some years since some applications are expected in the automotive industry. Several kinds of TWIP (Twinning Induced Plasticity) steels are now intensively studied. The aim of the present work is to elucidate the role of slip, twinning, and their mutual interactions on the stress-strain response of Fe-Mn-C TWIP steels thanks to a multiscale characterisation methodology. An initially fully austenitic grade (Fe – 22 wt % Mn – 1.2 wt % C) was tested in uniaxial tension up to necking. Light microscopy, EBSD and TEM were then used to characterise specimens strained to increasing levels of plastic deformation. In the early stage of plasticity (ε=0.02), the majority of the dislocations are perfect dislocations with a Burgers vector b=a/2<110>, and lying in the <111> plane. Dislocations are straight with pronounced edge character (Figure 1). Planar defects were also observed and identified as large overlapping stacking faults (OSF) (Figure 1). These defects can be considered as twinning precursor. The dislocations bounding the OSF were found to be Shockley PD with b=a/6<121> Burgers vector type, and the formation mechanism of deformation twinning could be accounted for by the three-layer twin model proposed by Mahajan and Chin [2]. After 10% of plastic strain, the microstructure is characterized by a pronounced deformation twinning, and dislocation cell formation between twin lamella. Simulation of the experimental SAD pattern shows that the crystallographic elements for twinning are K1={111} (primary twinning plane) and η1=<121> (primary twinning direction). Thanks to high resolution EBSD mapping, it is found that twinning occurs inhomogeneously between the different grains (Figure 2). A particular procedure analysis is under investigation in order to estimate the twinning rate as a function of the level of plastic deformation.
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1. 2. 3.
Nishiyama, Z., Oka, M. and Nakagawa, H., Trans. Japan Inst. Metals, 6, 1965, p. 88. Mahajan, S. and Chin, G. Y., Acta metal, 21, 1973, p. 1353. This work has been carried out in the framework of the IAP programme P6/24, “Physics based multilevel mechanics of metals”.
Figure 1. Two beam micrograph with the (1-11) reflecting plane, showing perfect dislocations and overlapping stacking faults
Figure 2. Crystallographic contrast map obtained by EBSD indicating twinning in some grains
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Misfit analysis of the InN/GaN interface through HRTEM image simulations J. Kioseoglou, G.P. Dimitrakopulos, Th. Kehagias, E. Kalessaki, Ph. Komninou, and Th. Karakostas Department of Physics, Aristotle University of Thessaloniki, GR-54124 Thessaloniki, Greece [email protected] Keywords: misfit dislocations, nitride interfaces, HRTEM simulations
Epitaxial growth of III-nitride compound semiconductors is based on specific interfacial structures, which affect the material properties and are defined by the structural mismatch between the abutting crystals. The (0001) wurtzite InN/GaN interface is characterized by a -10% in-plane mismatch between the lattices periodicities [1]. The high structural misfit leads to plastic relaxation of the strain that can be perceived by the presence of regularly spaced extra half-planes at the interface. These can be ascribed to a network of a-type (b=1/3< 2 1 10 >) misfit dislocations. The structure and energy of all admissible interfacial structures have been studied previously by molecular dynamics (MD) using the Tersoff empirical interatomic potential and a total of ~15,000 atoms [2]. In the present work, we undertake simulations of plan-view high resolution transmission electron microscopy (HRTEM) images, in order to elucidate the in-plane configuration of misfit dislocations observed in such interfaces. In Figure 1(a) a [0001] projection of the relaxed InN/GaN atomic interfacial configuration is shown, obtained by MD calculations. The corresponding HRTEM image simulations were performed using the EMS software package [3], resulting in maps of through-focus/thickness images for thicknesses varying from 1.7 up to 9.4 nm. The InN and GaN structures were almost equally sized through the thickness variation of the InN/GaN supercell. The HRTEM simulations resulted in configurations of the moiré fringes that are obtainable from the 0 110 reflections under the electron microscope Jeol 2011, with a point resolution of 0.19 nm and Cs=0.5 mm. It was evidenced that the local arrangement of the moiré fringes changed with increasing InN/GaN supercell thickness. Below ~5 nm supercell thickness (Figure 1(b-c)) the moiré fringes lie parallel to the { 10 10 } planes of both crystals as expected, but they intersected in pairs giving the impression of a “Star of David” network [4]. For a supercell thickness greater than ~5 nm (Figure 1(de)) the moiré fringes lie, locally, parallel to the { 1120 } planes forming zig-zagged configurations. Moreover, the latter intersected in triads forming a network of hexagons. It is concluded that the plan-view analysis of bad fit regions in high-misfit epitaxial bicrystals, based on moiré fringes, is dependent on the thickness of the sample and therefore, might prove tentative to the direct interpretation of the properties of the interfacial plastic relaxation. Nevertheless, multiple descriptions of high-misfit interfaces in
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terms of networks of misfit (primary) dislocations are feasible with all of them being consistent with the same physical bicrystal [5]. 1. 2. 3. 4. 5.
1. Th. Kehagias, A. Delimitis, Ph. Komninou, E. Iliopoulos, E. Dimakis, A. Georgakilas, Appl. Phys. Lett. 86 (2005), art. no. 151905. 2. J. Kioseoglou, E. Kalessaki, G. P. Dimitrakopulos, Ph. Komninou, and Th. Karakostas, J. Mater. Sci. (2007) (in press) DOI:10.1007/s10853-007-2235-0. 3. P. A. Stadelmann, Ultramicroscopy 21 (1987), p. 131. 4. A. M. Sánchez, J. G. Lozano, R. García, M. Herrera, S. Ruffenach, O. Briot, and D. González Adv. Funct. Mater. 17 (2007), p. 2588. 5. Marie Curie R T N “PARSEM” (MRTN-CT-2004-005583) is acknowledged.
Figure 1. (a) [0001] projection of the relaxed InN/GaN interface. (b)-(e) HRTEM simulated images of the InN/GaN supercell corresponding to the 1st and 2nd contrast maxima, for 1.7 nm thickness in (b)-(c) and 9.4 nm thickness in (d)-(e), respectively.
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Application of TEM for Real Structure Determination of Rare Earth Metal Compounds L. Kienle1, V. Duppel1, Hj. Mattausch1, M.C. Schaloske1, and A. Simon1 1. Max Planck Institute for Solid State Research, Heisenbergstr. 1, 70569 Stuttgart [email protected] Keywords: HRTEM, Electron diffraction, X-ray diffraction
Transmission electron microscopy (TEM) techniques are well adapted for the analysis of disordered crystals due to the minimization of space averaging effects. Hence, TEM enables the discovery of frequently overlooked, but new and chemically highly relevant aspects of solids, e. g. in the fields of rare earth metal (Ln) compounds. Samples with the nominal composition of La14(C2)7Cl9 contain crystals with prominent structural disorder. As a rule, the precession electron diffraction (PED) patterns exhibit faint diffuse streaks indicating a one-dimensional disorder. A first model of the real structure is proposed which is based on thick lamellae of a yet unobserved structural variant with the composition La14(C2)7Cl9. PED patterns without any diffuse scattering correspond with the metrics and calculated intensity of the new structural variant, cf. experimental (a) and simulated (b) PED pattern in Figure 1. Continuing X-ray analyses surprisingly demonstrate the existence of a homologous series of related, but ordered phases with unprecedented structural complexity [1], [2]. Carbide nitrides Ln4CNI6 (Ln = La, Ce, Gd) and Ln6(C2)NI9 (Ln = Y, Ce) are characterized by chains of octahedral and double-tetrahedral Ln6 building units which are centered by C2 dumbbells and N atoms, respectively [3]. A rotational disorder of the chains and variable sequences of the building units inside the chains give rise to a plethora of strongly disordered structures. All of them lead to characteristic anomalies in electron diffraction patterns, cf. cut-outs in Figure 2a. The patterns on the left and center correspond to a rotational disorder of every second and fourth chain, respectively. The diffuse double spots seen in pattern 2a (right) indicate broad modulations of the structure which are interrelated with a variable sequence of the building units inside the chains. The high resolution micrograph in Figure 2b was recorded on a selected area with a new type of complex ordering of the chains which are generally composed by alternating pairs of Y6 octahedra and double tetrahedra. 1. 2. 3.
Hj. Mattausch, et al. in preparation. Hj. Mattausch, C. Hoch, C. Zheng, A. Simon, Z. anorg. allg. Chem. 2007, 633, 239. Hj. Mattausch, H. Borrmann, R. Eger, R.K. Kremer, A. Simon, Z. anorg. allg. Chem. 1994, 620, 1889.
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Figure 1. Experimental (a) and simulated (b) PED pattern of a new structural variant of La14(C2)7Cl9 (zone axis [100], t = 10 nm).
Figure 2. a) Cut-outs of SAED patterns recorded on crystals with the nominal composition Y6(C2)NI9, see text, b) HRTEM micrograph and inserted simulation (zone axis [100], Δf = 25 nm, t = 4.5 nm).
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Quantitative Dislocation Analysis of 2H AlN:Si grown on (0001) Sapphire Oliver Klein1, Johannes Biskupek1, Ute Kaiser1, Sarad Bahadur Thapa2 and Ferdinand Scholz2 1. Central Facility of Electron Microscopy, University of Ulm, Germany 2. Institute of Optoelectronics, University of Ulm, Germany [email protected] Keywords: Aluminium Nitride, Sapphire, Dislocations
In the last few years aluminium nitride (AlN) has attracted much attention due to its extremely large direct band gap of approximately 6.0 eV and its impressive chemical and thermal stability. Thus AlN and AlxGa1-xN ternary alloys are promising materials for high-power high temperature electronic applications and optoelectronic devices in UV range. For group-III nitride wafers are still not available in sufficient amount and quality, AlN has to be grown on foreign substrates such as Al2O3 (Sapphire). Due to the lattice mismatch between the AlN/Al2O3 interface of about -11.7%, compressive stress is induced in the crystal system. The strain energy is reduced by the formation of threading dislocations, decreasing the crystal quality [1, 2]. Thus it is still a big challenge to grow AlN directly on foreign substrates with small dislocation density. To make the material suitable for semiconductor devices an efficient doping is necessary to achieve sufficient conductivity. Unfortunately Si doped AlN is still highly resistive mainly due to its large activation energy of several 100meV for the Si dopants. Thus high doping densities of up to 1020cm-3 are necessary, however affecting crystal quality. In this work we analysed the threading dislocations (TDs) quantitatively in highly doped ALN:Si layers by exploiting the 3g weak beam dark field method (WBDF) in cross-sectional transmission electron microscopy. The burgersvectors were determined by using the |g·b| criterion. To investigate the dislocations depending on the doping density, different doped 300nm thick AlN layers were grown on undoped ALN layers under same growth conditions by MOVPE (see [1, 3, 4]). It is shown that most dislocations formed in the undoped AlN layer are pure edge dislocations of type 1/3[21-10] along c axis, whereas the number of [0001] pure screw and 1/3[2-1-13] mixed dislocations is very small (Figure 1). Plane-view TEM investigations revealed a dislocation density of 4.1·1010cm-2 for the undoped AlN layer. When growing doped AlN:Si on the undoped layer under same growth conditions the pure edge dislocations penetrate the doped AlN layer without significant changing for low doping densities (doping density 1.5·1018cm-3). By increasing the doping density to 3.0·1019cm-3 the pure edge dislocations of type 1/3[2-1-10] in the undoped AlN layer change direction at the AlN/AlN:Si interface and band together (Figure 2). This annihilation of the pure edge dislocations is probably promoted by the increasing compressive stress resulting from the lattice mismatch between AlN and AlN:Si. It was also observed for highly doped AlN, that the propagation of the pure screw and mixed dislocations of type [0001] and
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1/3[2-1-13] is blocked at or near the AlN/AlN:Si interface by forming dislocation loops. This effects decrease the dislocation density at the surface. Beside the annihilation effect, the aggregation of the edge dislocations near the surface leads to a degradation of the surface quality (roughness) of the AlN:Si layer. 1. 2. 3. 4.
Thapa, S.B., et al., Structural and spectroscopic properties of AlN layers grown by MOVPE. Journal of Crystal Growth, 2007. 298: p. 383. N. Kuwano, et al., TEM analysis of threading dislocations in crack-free AlxGa1-xN grown on an AlN(0001) template. physica status solidi (c), 2003. 0(7): p. 2444-2447. Hertkorn, J., et al., Optimization of nucleation and buffer layer growth for improved GaN quality. Journal of Crystal Growth, 2007. 308(1): p. 30. Thapa, S.B., et al., MOVPE Growth of High Quality AlN Layers and Effects of Si doping. physica status solidi (a), 2008.
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b
Figure 1: WBDF images of the undoped AlN layer on (0001) sapphire from same sample area by using 0002 and 2-1-10 reflection: (a) Only TDs with b [0001] & 1/3[2-113] are visible, whereas the TD “1” is a pure screw and “2” a mixed dislocation. (b) Additional TDs are visible which are all pure edge dislocations of type 1/3[2-1-10]. a
b
Figure 2: Doped AlN:Si layer on undoped AlN (same growth conditions, doping density 3.0*1019cm-3): (a) WBDF image with g = (0002) shows pure screw and mixed dislocations forming dislocation loops at the AlN/AlN:Si interface. (b) Annihilation of the pure edge dislocations at the AlN/AlN:Si interface results in a decrease of the dislocation density at the surface.
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Transrotational crystals in crystallizing amorphous films: new solid state order or novel extended imperfection V.Yu. Kolosov Physics Dept. Ural State Economic University, 8Marta 62, Ekaterinburg, 620219 Russia [email protected] Keywords: extended crystal imperfection, amorphous-crystalline transformations, in situ TEM
As was first discovered for Se [1] and later on proved for some different substances amorphous-crystalline transformation in thin films can be associated with an unusual phenomenon: strong (up to 100-200 degrees per micrometer) regular dislocation independent lattice bending round an axis (or axes) lying in the film plane of the growing crystal. Other unusual (for condensed matter physics) crystal-like aggregation of atoms have been reported and widely recognized by the community in recent 20 years: quasi-crystals [2], fullerenes [3] and nanotubes [4] (followed by various nano derivatives). Anyway much less known crystals and structures with internal lattice bending (new term – “transrotation” [5]) during this period have been eventually recognized/studied in a large variety of thin film systems including well-known chalcogenide compositions [6] used for optical memory (i.e. rewritable formats of CD and DVD disks) and other non-volatile memory devices. Thus we suppose transrotational crystals nowadays deserve much more attention and discussion. In the present paper the main findings accumulated during the last period first and foremost by transmission electron microscopy (bend-contour technique [7]) are presented. High resolution electron microscopy and atomic force microscopy are used in due case to illustrate some details of both initial amorphous and crystallized areas and also their interface. Full set of the arguments and particulars available are presented to show how one can choose unambiguously between film bending/buckling and internal lattice bending (i.e. transrotation, that is translation of the unit cell accompanied by regular small rotation, Figure 1). In situ studies are presented for direct evidence of nucleation and growth of transrotational crystals in amorphous film under influence of the electron beam on TEM grid with a free standing amorphous film. These include multiple reversible local transformations "amorphous - crystalline (with transrotational microstructure)" recorded also as video for Se-Te films. Our TEM studies of transrotational crystals of different chemical nature, preparation and crystallization conditions are illustrated for Se-C, Se-Te, Sb2Se3, Sb2S3, Ge-Sb2Se3, Ge-Te, Tl-Se, Cu-Te, α-Fe2O3, Cr2O3, Co-Pd, Re, W, PZT ferroelectrics. Several observed types of internal geometry of transrotational crystals (including single crystals, whiskers - Figure 2, grains) are identified and described. They are the basis for description of different textures and orientational gradients of various microstructures and nanostructures observed in particular in polycrystalline films.
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Main factors (starting from film thickness and composition) that influence the degree and geometry of transrotation are presented and discussed. Atomistic mathematical model for the atom positions in "transrotational" single crystal corresponding to the conformal (or quasi-conformal) transformation of usual crystal lattice to the transrotational lattice is outlined, as well as hypothetical mechanism of unusual crystal growth in an amorphous matrix based on the surface nucleation. Generally transrotational crystals (structures) revealed by TEM can be either considered as a new state intermediate between glassy and crystalline states (similarly to the structure of liquid crystals regarded as intermediate between crystalline and liquid), or regarded as an example of new kind of defects (truly extended) in condensed matter: 3D imperfection that can named as “transrotations” which supplements well known dislocations (in crystals) and disclinations (in liquid crystals). 1. 2. 3. 4. 5. 6. 7.
I.E. Bolotov, V.Yu. Kolosov and A.V. Kozhyn, Phys. Stat. Sol., 72a (1982), p.645. D. Shechtman et. al. Phys. Rev. Let., 53 (1984), p.1951 H. Kroto et. al., Nature, 318 (1985), p. 162. I. Iijima, Nature, 354 (1991), p.56. V.Yu. Kolosov and A.R. Tholen, Acta Mat. 48 (2000) p. 1829. B.J. Kooi and J.Th.M. De Hosson, J. Applied Physics 95 (2004), p. 4714. V.Yu. Kolosov, Proc. XII ICEM, ed. L. Peachey & D. Williams, (Seattle) 1 (1990) p. 574
a b Figure 1. Schematic cross section through a thin film crystallized from an amorphous film, showing the way in which the orientation of the crystal planes changes continuously from place to place - (a), deduced from micrographs similar to one presented in Figure 2. Two sets of crystal planes {100} and {010} are sketched (approximately 1 line per 10 crystal planes) – (b).
Figure 2. A micrograph of a grain in a crystallized region of a film of Cu-Se. The bend contours crossing the grain, along its length, indicate the change of orientation, which is of the kind shown schematically in Figure 1.
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Determination of precise orientation relationships between surface precipitates and matrix in a duplex stainless steel Y. Meng1,2, G. Nolze3, W.Z. Zhang1, L. Gu2, P.A. van Aken2 1. Laboratory of Advanced Materials, Department of Materials Science and Engineering, Tsinghua University, Beijing, China. 2. Max Planck Institute for Metals Research, Heisenbergstr. 3, 70569 Stuttgart, Germany 3. Federal Institute for Materials Research and Testing (BAM), Berlin, Germany [email protected] Keywords: Orientation relationship, surface precipitates, precipitation crystallography, duplex stainless steel
The knowledge of the strain field due to precipitation is essential for understanding the evolution of the microstructure during phase transformation. The transformation strain may be partially relaxed, but quantitative description of the residual elastic strain is an unsolved problem. The body-centre-cubic (bcc) structured ferrite (δ) to facecentre-cubic (fcc) structured austenite (γ) phase transformation in bulk material in a duplex stainless steel presents a rather consistent orientation relationship (OR) [1], but a preliminary investigation on the precipitates near the surface reveals the OR different from that in the bulk [2]. The constraint to the precipitates transformed in the near surface layer is different from those in bulk, so crystallographic feathers are expected different. A systematic and precise measurement of the OR of the precipitates in the near surface layer will provide useful experimental data for examining and modelling the effect of local constraint interacted with the strain field generated from precipitation in development of the OR and the precipitate morphology. In the present work the OR between δ and γ phases has been precisely measured with both Kikuchi electron diffraction pattern and electron back-scattered diffraction pattern (EBSP) techniques. Kikuchi patterns obtained from different phases can be used to define uniquely a set of three parallel directions, sufficient for defining the OR with satisfied accuracy [3]. Fig. 1(a) and (b) are a pair of Kikuchi patterns from γ and δ phases. Usually at least 3 sets of measurements are made from one site to reduce experimental error and uncertainty in the variants selection. The analysis of Kikuchi patterns is assisted by plotting stereographic projections for a systematic analysis; Fig. 1(c) and (d) contain three pairs of Kikuchi pattern measurements from fcc and bcc. The determined OR is expressed to two ways: one in form of an OR matrix for further calculation and the other in the convention of the angles between conjugate closepacked planes and directions. In addition to the precise but local TEM measurements, Euler angle maps are obtained by the EBSP technique for examining the OR in a broad surface area [4]. A set of Euler angles specify the spatial orientation of a crystal. Therefore the relative Euler angles along an interface define the OR (Fig. 2(a) and 2(d)). Two types of irrational ORs have been determined. One is near the NishiyamaWassermann (NW) OR ((1 1 1)f||(101)b, [110]f||[010]b, where the subscripts f and b stand S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 659–660, DOI: 10.1007/978-3-540-85226-1_330, © Springer-Verlag Berlin Heidelberg 2008
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for fcc and bcc lattices respectively). This OR can be expressed as: [110]f||[ 0.028 1 0.033]b, 2.5˚ from [010]b, (111) f || (1 0.005 0.975)b , 0.75˚ from (101)b. The precipitates
having this OR are in contact with (and possibly nucleated from) a δ grain boundary and hold this OR with respect to the matrix grain into which the precipitates grow preferentially, as shown in Figure 2(a). They have a broad and straight interface as shown in the TEM image in Figure 2(b). The other type of ORs is near the Kurdjumov– Sachs (KS) OR ((111)f||(011)b, [01 1 ]f||[ 1 1 1 ]b): [01 1 ]f||[ 0.993 0.976 1 ]b, 0.48˚ from[ 1 1 1 ]b, (111)f || (0.061 1 0.946)b, 3.0˚ from (011)b. The precipitates having this OR are found in a matrix grain and they exhibit an acicular shape (Fig. 2(c)). Some of them are present in pairs, holding a twinning relationship with each other (Fig. 2(d)). Further experimental and theoretical studies for understanding the observations are undergoing. 1. 2. 3. 4. 5.
D. Qiu and W.Z. Zhang, Acta Materialia, 55 (2007), p.6754 Qiu and W.Z. Zhang, Acta Metal. Sinica vol. 41 (2005). p897 M.-X. Zhang and P.M. Kelly, Scripta Materialia, Volume 37(1997), Number 12, p. 2009 G. Nolze and V. Geist, Cryst. Res. Technol. 39(2004), No. 4, 343 The support of National Natural Science Foundation of China Grant No. 50471012 and 50671051 and aid of the Ph.D. Foundation of Education Ministry are acknowledged
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d
c
b
Figure 1. Kikuchi pattern from the grain boundary precipitate and SAD pattern from its matrix: (a) [1 0.430 0.247]f and (b) [131]b. The stereographic projections of fcc (c) and bcc (d) near the measurements, in addition to the data in (a) and (b) (indicated by arrows), two pairs ([0.975 0.214 0.065]f||[351]b and[0.946 0.304 0.109 ]f||[120]b) were also plotted as on the other corners of the triangles, used to verify the OR.
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b
c
=2 µm; Map6; Step=0.3 µm; Grid50x50
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Figure 2. The precipitates along a matrix grain boundary grow preferentially towards one side. The interfaces are indicated by yellow lines in the EBSP map (a) that stand for a deviation from NW OR less than 3º. The TEM image (b) shows the precipitate has a wide straight interface. The precipitates distribute in a single matrix grain (c). The green line in (d) is a detected Σ3 interface in Euler angle map along a precipitate’s main axis.
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Interfaces in Cu(In,Ga)Se2 thin film solar cells G. Östberg, E. Olsson Chalmers University of Technology, Department of Applied Physics, SE-412 96 Göteborg, Sweden [email protected] Keywords: TEM, interfaces, CIGS
The demand for solar cells as energy harvesting, increases all the time. Cu(In,Ga)Se2 (CIGS) is a strong candidate for use as p-type absorption layers in thin film solar cells. Such cells are built up by a number of thin layers applied on a supporting substrate, usually soda lime glass (SLG). As the cell is irradiated, photons are absorbed in the CIGS layer and electron/hole pairs are created, which causes a current to flow. Since the charges are transported across the layers they will pass through a number of interfaces. Dangling bonds and other defects at the interfaces will act like recombination centres where the charges will be trapped. Therefore, the geometry and chemistry of these interfaces are very important parameters for the performance of the solar cell. In this study, interfaces between some different layers in a CIGS thin film solar cell have been investigated by transmission electron microscopy (TEM), high resolution TEM (HRTEM), energy filtered TEM (EFTEM) and high angle annular dark field (HAADF) analysis. The cells were prepared by depositing a 1.5 µm CIGS absorbing layer mainly by sputtering on a SLG substrate covered with a 400 nm layer of Mo, used as back contact. On top of the CIGS, a 50 nm CdS buffer layer was applied by chemical bath deposition (CBD) and a 400 nm ZnO layer was sputtered. Finally, a transparent top contact of InSnO2 (ITO) was sputter deposited. A TEM cross-section of the complete cell is shown in Figure 1. Thin foil TEM samples were prepared by in situ lift-out in an FEI FIB-SEM workstation and were attached to a supporting Ti grid. In order to reduce the surface damage and Ga implantation from the FIB-SEM preparation, the foil was ion polished for 60 seconds in a Gatan PIPS system using ±10º beam angle at 2.0 keV. In Figure 2a a close-up of the interfaces on top of the CIGS is shown. The CIGS/CdS and ZnO/ITO interfaces are very smooth and there are no obvious signs of defects. However, the CdS appears to have been damaged by the sputtering of the ZnO and voids have formed between the two layers. Obviously, these voids will hamper the charge transport and thereby increase the series resistance of the cell. Figure 2b shows a close-up of the CIGS/Mo interface. The Mo layer has a clear columnar grain structure with grains about 50 nm in width whereas the CIGS grains are much wider with sizes up to 0.5 µm. The CIGS layer is very well attached to the Mo layer and no voids can be seen at the interface.
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ITO ZnO CdS
CIGS
Mo Figure 1. TEM cross-section of the studied solar cell.
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b
Figure 2. Close-ups of the interfaces between CIGS, CdS, ZnO, and ITO (a) and between CIGS and Mo (b).
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In-situ electron beam irradiation of nanopipes in GaN F. Pailloux and J.-F. Barbot Laboratoire de Métallurgie Physique, UMR 6630 CNRS/Université de Poitiers, Bâtiment SP2MI, Téléport 2, BP30179, 86962 Chasseneuil Futuroscope Cedex, France [email protected]
Keywords: gallium nitride, dislocation, electron irradiation damage Gallium nitride (GaN) is one of the promising wide band gap (Eg=3.39 eV) semiconductor for optoelectronic applications. When epitaxially grown on bufferedsapphire substrate, GaN films often exhibit a rather huge density of structural defects such as threading dislocations (up to109 cm-2) and nanopipes (up to 106 cm-2). A lot of work is devoted to the understanding of the structure and the nucleation of these defects and of their influence on the optoelectronic properties of GaN. Among the scenarii suggested to explain the formation of open core dislocations, the segregation of doping species or impurities (especially oxygen) to the dislocation core seems well admitted. In a recent study, Hawkridge & Cherns [1] have pointed out that nitrogen can be substituted by oxygen at surface pits during the growth of the film. By lowering the surface energy of {10-10}-planes, oxygen impurities would thus promote the opening of the dislocation core (only for screw-type dislocations) and would consequently lead to the growth of a nanopipe. It has also been reported that some screw dislocations have alternate open and closed core segments, which would suggest that the core structure of dislocation (or the nanopipe morphology) is driven by oxygen impurities concentration. Nevertheless, in a previous paper [2], we have shown that regular nanopipes experience a morphological change (namely a “pinch-off “) under 300kV-electron irradiation (in a JEOL 3010 TEM, LaB6, 0.19nm resolution), leading to chains of pinholes linked by a screw dislocation (Figure 1). We described this transition in terms of surface energy lowering and suggested that the starting point was the nucleation of point defects (interstitial and vacancies) by the electron beam, which leads to a roughening of the lateral surface of the nanopipe. We argued that, under the irradiation process, the point defects located close to a nanopipe, which initial facets are of {10-10}-type, promote the surface relaxation to form pinholes with {10-11}-type facets that are suspected to be energetically more stable than the {10-10} ones (as they are similar to {111}-facets in the sphalerite structure). We have carried out new detailed investigations of such morphological changes. Despite the role of oxygen in this transition is still unclear, we give a precise drawing of the kinetics of this transformation. We point out that such a final configuration (chain of pinholes linked by a c-screw dislocation) is more stable than a single open core dislocation (nanopipe) and would like to temper the role of oxygen or other impurities in the nucleation of such defects. 1. 2.
Hawlridge M.E. and Cherns, D., Appl. Phys. Lett. 87 (2005) 221903 Pailloux F., Colin J., Barbot, J.F. and Grilhé J., Appl. Phys. Lett. 86 (2005) 131908
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Figure 1. Transition of a regular nanopipe in a chain of pinholes. (a) Regular nanopipe (starting of the experiment). (b) During the first step of the transition, point defects are created in GaN; at the same time, roughness appears on the lateral surfaces of the nanopipe. (c) After long time irradiation, the nanopipe is entirely dissolved as a chain of pinholes.
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The atomic structure of an incommensurate (001)-(110) Si grain boundary resolved thanks to a probe Cs-corrector J.L. Rouviere1, F. Lançon2, K. Rousseau1,*, D. Caliste2 and F. Fournel3 CEA-Grenoble 17 rue des Martyrs 38054 Grenoble FRANCE 1.INAC/SP2M/LEMMA 2. INAC/SP2M/L_SIM 3. LETI –Minatec * Now at Serma Technologies BHT-Bat 52 7 Parvis Louis Néel, BP 50 38040 Grenoble France [email protected] Keywords: Z-contrast, STEM, probe Cs-corrector, grain-boundary, incommensurate
A non-symmetric and non-periodic (001)/(110) Si grain boundary (GB) has been experimental studied by transmission electron microscopy (TEM) and its atomic structure refined by atomic simulation. The (001)/(110) Si grain boundary was made by bonding together two different 4” Si wafers, having respectively (001) and (110) surfaces. The wafers were carefully aligned in order that they have a common in-plane [110] direction. In the interface perpendicular to [110], the [110] direction of the first grain is then parallel to the [001] direction of the second grain (figure 2). As the ratio of the norms of these two directions is equal to 2 , it is mathematically impossible to have a common period along these directions and the GB is said to be incommensurate [1]. For TEM observations, cross-section samples were classically prepared by mechanical polishing followed by ion milling. High Resolution (HR) TEM images were first realised on a JEOL 400KV microscope with a LaB6 filament. These images (like figure 1) give a general view of the GB. The interface is not flat but has a zigzag shape. The interface is not periodic, but composed of a limited number of segments distributed at first glance with no apparent order. These segments are characterised by their respectively lengths p1=k1/2[110]1 and p2=k2[001]2 in the lower and upper crystals (convention of figure 2). Figure 1 shows the more frequent (k1,k2)-segments : (7,5), (10,7), (13,9) and (3,2). It can be checked that the k1/k2 ratios are good approximants to 2 although they do not all belong to the diophantian series [1]: (1,1), (3,2), (7,5), (17,12). This and the order of (k1,k2)-segments will be discussed in the presentation. It was impossible to solve the atomic structure of these segments by analysing HR-TEM images because the structures are too complex. We had to use Z-contrast images to do it. HAADF Z-contrast STEM images were realised on a FEI TITAN microscope working at 300kV and equipped with a CEOS probe Cs-corrector. Figure 2 was obtained with a spot size of 6, a gun lens of 6 an extraction voltage of 4300V, a convergent angle of 20mrad, a collection angle of 40mrad and a scanning time of 20s. The original image has been unsharped (smoothed) over 7x7pixels regions to enhance the contrast. The atomic structure can be directly deduced from figure 2: each white spot corresponds to an atomic column and Si dumbbells composed of two columns separated by 0.135 nm are resolved. Several isolated columns can be seen and correspond to a special reconstruction doubling the period along [110]. The atomic structures of several (k1,k2) segS. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 665–666, DOI: 10.1007/978-3-540-85226-1_333, © Springer-Verlag Berlin Heidelberg 2008
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ments and their energies computed using by a Stillinger-Weber potential (figure 3) will be discussed. Models were built with the help of the V_Sim software [2]. 1 2
F. Lancon, J.M. Penisson, U. Damen Europhys. Lett., 49 (2000), p. 603 D. Caliste, http://inac.cea.fr/L_Sim/V_Sim
Figure 1 High-resolution TEM image of the incommensurate (001)/(110) GB. The crystals are relatively thick and not perfectly oriented along the common [110] direction (which is perpendicular to the paper) in order that the grains have different contrasts: the zig-zag shape of the interface is then outlined. Numbers above and below the interface give the (k1,k2) numbers associated with segment lengths.
Figure 2 High Resolution STEM image showing several periods of the (7,5) segment. The common [011] direction of the two grains is the direction of observation. Arrows points towards atomic columns that have special reconstructions (no dumbbell) doubling the period along [011]. The white squares outline specific tunnels composed of rings of 7 atoms. (see figure 3)
Figure 3 Atomic model of the (7,5) segment of the incommensurate GB projected along [011]. It can be directly compared to figure 2. Arrows and rectangle correspond to those of figure 2
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Atomic structure and dopant segregation of [0001] tilt grain boundaries in ZnO bicrystals Y. Sato1, T. Mizoguchi2, J.P. Buban2, N. Shibata2, T. Yamamoto3, T. Hirayama1, and Y. Ikuhara1,2 1. Nanostructures Research Lab., Japan Fine Ceramics Center, Nagoya, Japan 2. Institute of Engineering Innovation, University of Tokyo, Tokyo, Japan 3. Department of Advanced Materials, University of Tokyo, Kashiwa, Japan [email protected] Keywords: grain boundary, segregation, ZnO, HRTEM, HAADF-STEM
Physical properties of ceramics are often changed by presence of grain boundaries (GBs). One possible cause of the change is the grain boundary (GB) itself, because the GB atomic structure is different from that of the bulk. Another possible cause would be dopant segregation at the GBs. Thus, it is important to understand GB atomic structure and dopant segregation at the GBs at the atomic level. Here, we are going to present our studies on atomic structure and dopant (Pr) segregation at ZnO [0001] tilt GBs. We fabricated undoped and Pr-doped [0001] tilt GBs within ZnO bicrystals [1]. Atomic structure of the undoped ZnO GB was observed by high-resolution transmission electron microscopy (HRTEM), and that of the Pr-doped ZnO GB was observed using high-angle annular dark-field scanning TEM (HAADF-STEM). Atomistic calculations were performed for both of the undoped and the Pr-doped ZnO GBs to obtain further information in detail. The calculation results were compared with the HRTEM and the HAADF-STEM images. Figure 1 shows a HRTEM image of the undoped ZnO Σ7 GB. It was found that undoped ZnO [0001] Σ7 GB has two kinds of stable structures. The two kinds of stable structures were characterized by structural units (SUs) as indicated in the figure. One SU has the characteristic atoms that have the coordination number of three, and the other has the characteristic atoms that have the coordination number of five. (Coordination number of ZnO bulk crystal is four.) It was found that undoped ZnO Σ49 GB has characteristic arrangement of the SUs along the GB plane. Two alternative Σ7like SUs (a and b in Fig. 1) and a bulk-like SU were arranged in this order. It is considered that the characteristic arrangement of the SUs effectively relaxes local strain at the Σ49 GB. Figure 2 shows a HAADF-STEM image of the Pr-doped ZnO Σ7 GB. It was found that two ZnO crystals were directly bonded at the GB. Brighter spots are found periodically along the GB. Since intensity of bright spots in HAADF-STEM image increases as atomic number included in the atomic columns increases, the brighter spots indicate the presence of atoms with higher atomic number, in this case, Pr. Comparison of the image with calculation results indicates that Pr substitutes specific Zn sites of the GB [2].
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Y. Sato, T. Yamamoto, and Y. Ikuhara, J. Am. Ceram. Soc., 90, (2007). p.337-357. Y. Sato, J. P. Buban, T. Mizoguchi, N. Shibata, M. Yodogawa, T. Yamamoto, and Y. Ikuhara, Phys. Rev. Lett., 97, (2006). p.106802-1-4.
Figure 1. HRTEM image of undoped ZnO [0001] Σ7 GB. Structural units (SUs) of the GB are shown by while quadrilaterals. Two kinds of SUs found are designated by “a” and “b”.
Figure 2. HAADF-STEM image of Pr-doped ZnO [0001] Σ7 GB. Bright spots in the image show the positions of atomic columns, which is different from in Fig. 1. Arrows indicate much brighter spots at the GB.
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TEM study of strain and defect engineering with diluted nitride semiconductors J. Schöne1,2, E. Spiecker1, F. Dimroth2, A.W. Bett2, W. Jäger1 1. Faculty of Engineering, Christian-Albrechts-University, Kaiserstrasse 2, D-24143 Kiel 2. Fraunhofer Institute for Solar Energy Systems ISE, Heidenhofstrasse 2, D-79110 Freiburg [email protected] Keywords: dislocation blocking, strain engineering, diluted nitrides
Optoelectronic semiconductor devices fabricated by hetero-epitaxy, such as heterostructure lasers and III-V multi-junction solar cells, need high material qualities and low dislocation densities. We have developed a technique controlling strain relaxation and dislocation population in metamorphic semiconductor heterostructures using diluted nitride intermediate layers [1]. Diluted nitride semiconductors, like GaAs1-xNx and GaP1-xNx, are known to be significantly harder than the corresponding nitrogen-free binary alloys, GaAs and GaP [2]. Possible factors contributing to the hardening via reduction of dislocation mobility or pinning of dislocations are the presence of nitrogen interstitials [3], the increased bond strengths associated with substitutional nitrogen [3] and local strain due to the small covalent radius of substitutional nitrogen [2]. By introducing diluted nitride intermediate layers with appropriate nitrogen concentrations into step-graded buffer structures we could demonstrate that diluted nitride layers can be used as “dislocation blockers” which effectively suppress misfit dislocations in predefined parts of the structures despite of the accumulation of strain [1]. Figures. 1 and 2 summarize the main results of our study by comparison of three different buffer structures. All structure were fabricated by metal organic vapour phase epitaxy (MOVPE) in an Aixtron AIX2600-G3 reactor on vicinal GaAs (001) substrates with 6° miscut toward (111)A. The TEM bright field image of a step-graded GaAs1-xPx/GaAs buffer without diluted nitride layers (Figure 1, left) reveals a homogeneous and continuous dislocation network throughout the graded part of the structure. In contrast, a step-graded GaAs1-xNx/GaAs buffer with similar misfit profile (Figure 1, right) shows dislocations only in the bottom part of the structure containing layers with nitrogen concentrations < 2 %. We conclude that x ~ 2 % represents the threshold nitrogen concentration for efficient dislocation suppression in GaAs1-xNx. Figure 2 (left) depicts a TEM bright-field image of a stepgraded GaAs1-xPx/GaAs buffer structure in which two intermediate GaAs1-xNx layers with x ~ 2.5 % and x ~ 4 % have been introduced. In accordance with the observed threshold nitrogen concentration misfit dislocations are effectively blocked already by the first GaAs1-xNx layer with x ~ 2.5 %. The absence of misfit dislocations in the layers above resulted in accumulation of tensile strain eventually leading to film cracking. This interpretation is corroborated by high-resolution XRD measurements. The scattering S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 669–670, DOI: 10.1007/978-3-540-85226-1_335, © Springer-Verlag Berlin Heidelberg 2008
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intensity in the asymmetric (2-24) reciprocal space map (Figure 2, right) indicates almost complete strain relaxation in the bottom part of the buffer and apparently no further strain relaxation in the top part. Our results demonstrate the potential of intermediate dilute GaAs1-xNx layers for controlling defect formation and strain relaxation during heteroepitaxial growth. 1. 2. 3.
J. Schöne, E. Spiecker, F. Dimroth, A.W Bett, W. Jäger, Appl. Phys. Lett. 92 (2008) 081905. K. Momose, H. Yonezu, Y. Fujimoto et al., Jpn. J. Appl. Phys., Part 1 41 (2002) 7301. M. Adamcyk, J. H. Schmid, T. Tiedje et al., Appl. Phys. Lett. 80 (2002) 4357.
Figure 1: TEM bright field image of the GaAs1-xPx/GaAs buffer structure showing a homogeneous dislocation network spreading out over the whole graded buffer (left). TEM dark field image of the GaAs1-xNx/GaAs buffer structure revealing an abrupt suppression of dislocation formation in the upper buffer layers (right).
Figure 2: TEM bright field image of a GaAs1-xPx/GaAs buffer structure with two intermediate GaAsN layers showing the dislocation blocking phenomenon (left). (2-24) RSM illustrating the corresponding effect on the layer strain (right).
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Investigation of the Co-Precipitation of Copper and Nickel in Silicon by Means of Transmission Electron Microscopy C. Rudolf1, L. Stolze and M. Seibt 1. IV. Physikalisches Institut, Georg-August-Universität Göttingen, Friedrich-Hund-Platz 1, 37077 Göttingen, Germany [email protected] Keywords: transition elements in silicon, co-precipitation, multi-metal silicides, FIB preparation, EDX,
We investigated the composition of multi-metal silicides formed by co-precipitation of copper and nickel in silicon by means of energy-dispersive x-ray analysis (EDX). Common multi-crystalline solar cell silicon is contaminated with a variety of transition metals among which iron, nickel and copper are the most prominent ones [1]. At room temperature the contaminants are predominantly precipitated. While the precipitation behaviour of single metal impurities in silicon is studied in detail during the past decades, the co-precipitation of different metals simultaneously present in silicon is an interesting new topic of investigation. Samples made by wafer-bonding of two single crystalline silicon wafers were codoped with both metals by thermal evaporation and diffusion annealing at 1050°C. After annealing the samples were cooled moderately fast with a rate of 6 K/s to enable formation and growth of metal precipitates on the grain boundary between the two wafers. Transmission electron microscopy (TEM) and energy-dispersive x-ray analysis reveals that two types of precipitates are formed as it is shown in Figure 1: copper rich silicide particles that contain a small amount of nickel in the order of 5 at.-% and nickel rich particles containing up to 20 at.-% of copper. These precipitates can be assigned to the known binary metal silicide phases Cu3Si and NiSi2 and a solid solution of a second metal species therein. No precipitates consisting of a ternary silicide were observed. TEM and EDX was performed on cross sectional samples. Regions of interest (ROI) that contain precipitates were first identified by defect etching as it is shown in fig 2A. Then the focussed ion beam technique was applied to prepare a sample from a ROI that is marked by a etch pit, compare Figure 2B. 1.
A.A. Istratov, T. Buonassisi, R.J. McDonald, A.R. Smith, R. Schindler, J.A. Rand, J.P. Kalejs, E.R. Weber, J. Appl. Phys. 94 (2003), p. 6552.
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A)
B)
1 2 Figure 1. A) High-resolution micrograph of a pair of a multi metal silicide particles whose compositions are shown by the EDX –spectra (B) to be either Ni-rich (at position 1) or Ni-rich (at position 2). A)
B)
ROI
Figure 2. A) Optical micrograph of a cross sectional sample after defect etching. A region of interest (ROI) is marked by a etch pit. By application of the focussed ion beam technique (b) a sample for the TEM analysis is prepared from the ROI.
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Micro-structure analysis of a friction-stir welded 2024 aluminium alloy using electron microscopy E. Sukedai1, T. Maebara1 and T. Yokayama1 1. Okayama University of Science, Okayama 700-0005, Japan [email protected] Keywords: friction-stir welding, 2024 Al alloy, helical dislocation, precipitation, EDS analysis.
Friction stir welding (FSW) method has been used in the wide field. In order to establish the reliability, micro-structures in the welded part have been investigated. The welded parts were investigated using electron microscopy. Some results obtained were different to the reported results. FSW was performed under a condition of tool tilt angle = 3 degrees, rotation speed = 1,300 rpm and welding speed = 330 mm/min. Thin foil specimens from the nugget and the mother material were prepared. Micro-strucutre observation and EDS analysis were carried out using a JEM 2010F electron microscope. Figure 1(a) and (b) show dislocation structures in the mother material and the nugget part, respectively. In Figure 1(a), lots of short dislocation segments are visible, while in Figure 1(b), helical dislocations indicated by an arrow and long segments are visible. Since helical dislocations are formed by vacancy concentration to dislocations, it is suggested that the friction-stir welded part was heated by the processing, and lots of vacancies were formed. Figure 2(a) and (b) show precipitations in the mother material and the nugget part, respectively. In Figure 2(a), rod-shape precipitations are aligned along a crystallographic orientation. Their size is 50-500 nm. In Figure 2(b), the rod-shape and round-shape precipitations are visible. The size of the former is approximately 200 nm. The density of precipitations in the nugget is lower than that in the mother material. This is considered as one of the reason of lower hardness value in the nugget area. Figure 3 shows an EDS spectrum of a rod-shape precipitation in the nugget, and Al, Cu, Mn, Cr, Si and Mg were detected. In the round-shape precipitation in the same area, Al, Cu, Fe and Mn were detected, but Si and Mg were hardly detected. In a literature [1], precipitations in the FSW parts of the same alloy are S-phase (Al2MgCu) and Ωphase (Al2Cu). EDS results of the present work indicate that the precipitations shown in Figure 2(b) are different to S-phase and Ω-phase. The diffraction pattern analysis is processing, and the results will be presented in the conference. EDS spectra of the matrix neighbouring the precipitations in the nugget indicates Al, Cu, and Mg. EDS analysis of rod-shape precipitations in the mother material was also carried out. The results were also different to the reported results. 1. M.J. Jones et al.: Scripta Materialia, 53 (2005) 693-697.
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b
Figure 1 Weak beam images of dislocation configurations in the mother material (a) and the nugget part in (b). Scale bar indicates 200 nm
a
b
Figure 2 Precipitations in mother material (a) and nugget part (b). Scale bar indicates 500 nm.
Figure 3 EDS spectrum from a rod-shape precipitation in the nugget part.
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Deformation defects in a metastable β titanium alloy H. Xing and J. Sun Electron Microscopy Laboratory, School of Materials Science and Engineering, Shanghai Jiaotong University, Shanghai 200240, P. R. China [email protected] Keyword: deformation defect, titanium alloy
The β or metastable β titanium alloys can be achieved by addition of d-electron rich elements, such as Nb, Mo, Ta and so on and an example is Ti-23Nb-0.7Ta-2Zr-1.2O (mol %). From the viewpoint of elastic stability of crystals, the critical valence electron number e/a associated with β to α transition for titanium alloys is about 4.24, and the alloy with this special value of e/a has a high elastic anisotropy, which results in a unique dislocation-free plastic deformation mechanism and mechanical property [1]. In this work, the deformation defect has been investigated by high-resolution transmission electron microscopy (HRTEM) in a metastable β titanium alloy with a chemical composition of Ti-23Nb-0.7Ta-2Zr-1.2O (mol %) after substantial cold swaging. This composition converts to an e/a of 4.24. The HRTEM observations showed that there are large accumulations of local strains and inhomogeneously lattice distortions in the alloy after substantial cold swaging. A few dislocations with {110} extra half planes can be clearly found from these figures and the Burgers circuits are drawn to enclose the cores of these dislocations, giving the projections of Burgers vectors of dislocations as 1/2<111>, as shown in Figure1. Additionally, the mechanical twins with crystalline orientation of <111>{112} were often observed in the deformed alloy, as shown in Figure 2. The HRTEM results indicates that this alloy plastically deforms via the conventional dislocation glide along <111>{110}, {112} or {123} and mechanical twin along <111>{112}, rather than the dislocation-free mechanism. The relationship between the deformation defect and the phase stability for the Ti-23Nb-0.7Ta-2Zr-1.2O alloy is finally discussed with above HRTEM observations. 1. 2.
T. Saito, T. Furuta, et al., Science 300 (2003), p. 464. We acknowledge the financial support from the National Natural Science Foundation of China (Grant No. 50571063) and the Science and Technology Committee of Shanghai Municipal (Grant No. 04JC14054).
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Figure 1 HRTEM images taken along [100] (a) and [111] (b) directions, respectively, showing dislocations in Ti-23Nb-0.7Ta-2Zr-1.2O alloy after substantial cold swaging. (b)
Figure 2 Conventional TEM (a) and HRTEM (b) images taken along [110] direction showing mechanical twins in Ti-23Nb-0.7Ta-2Zr-1.2O alloy after substantial cold swaging.
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Defect generation and characterization in 4H-SiC J.P. Ayoub, M. Texier, G. Regula, M. Lancin and B. Pichaud Aix-Marseille Université, IM2NP CNRS, IM2NP (UMR 6242) Faculté des Sciences et Techniques, Campus de Saint-Jérôme, Avenue Escadrille Normandie Niemen - Case 142, F-13397 Marseille Cedex, France [email protected] Keywords: stacking faults, dislocation core, SiC
Study of extended defects in 4H-SiC actually receives particular attention since high quality samples are now available while mechanisms of defect nucleation and propagation are still subject of debate. Indeed, several works revealed that Schockley partial dislocations easily propagate in the brittle regime, leading to the extension of simple [1] and multiple stacking faults [2,3] depending on the deformation conditions. Some of them also bring forth the influence of the n doping on the stacking fault multiplicity [4]. In addition, recent experiments suggest that Si(g) core segments are more mobile than C(g) dislocations [1,5] whereas opposite behaviour is expected from atomistic simulations [6]. Intensive works are thus now focused on the determination of the respective influences of doping and mechanical stress on the defect nature in 4HSiC. The aim of our fundamental study is to produce and propagate planar defects in SiC by surface scratching, cantilever bending and annealing, and to identify their multiplicity and the core composition of bounding dislocations. This paper deals with the defects which propagate in highly N-doped (11-20) oriented 4H-SiC wafers. Chemical etching of the surface reveals characteristic etch patterns corresponding to the emergence of planar defects at the surface. Few families are distinguished from the propagation direction and extension of the defects. An examination of the etch pattern tips by optical and atomic force microscopy also allows us to discriminate various facieses for each defect family. The complete characterization of stacking faults and partial dislocations is performed by weak-beam dark-field and high resolution imaging completed by large angle convergent beam electron diffraction (LACBED) analyses. The core composition of observed dislocation segments is deduced from dislocation core reconstruction. We discuss the possible relationships between the characteristics of the structural defects and the features of the corresponding etch patterns. The mobility of the bounding partial dislocations is derived from the stacking fault extension and correlated to their character and core composition. 1. 2. 3.
P. Pirouz, J.L. Demenet, M.H. Hong, Phil. Mag. A, 81, (2001), p. 1207. H.J. Chung, J.Q. Liu, M. Skowronski, Appl. Phys. Lett., 81 (2004), p. 3759. A. Mussi, J. Rabier, L. Thilly, J.L. Demenet, Phys. Stat. Sol., 4 (2007), p. 2929.
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4. 5. 6.
M. Zhang, H.M. Hobgood, M. Treu, P. Pirouz, Mat. Sci. Forum, 457-460 (2004), p. 759. M. Texier, G. Regula, M. Lancin, B. Pichaud, Phil. Mag. Lett., 86 (2006), p. 529. A.T. Blumenau, C.J. Fall, R. Jones, S. Öberg, T. Frauenheim, P.R. Bridon, Phys. Rev. B, 68 (2003), p. 174108.
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Investigation of defects in polymorph B enriched zeolite Beta Daliang Zhang1,2, Junliang Sun1,2, Sven Hovmöller1, Xiaodong Zou1,2* 1. Structural Chemistry, Stockholm University, SE-106 91 Stockholm, Sweden 2. Berzelii Centre EXSELENT on Porous Materials, Stockholm University, SE-106 91 Stockholm, Sweden [email protected] Keywords: zeolite Beta, HRTEM, defects
Defects occur in many crystals. Stacking fault is one kind of planar defects. In structures of porous materials, channels and cages can be considered as stacked according to certain stacking sequences. Disruption of these stacking sequences leads to stacking faults or twins. It is very important to study the defects in a material in order to understand the structure information and the properties of the material. High silica zeolite Beta is one of the industrially important zeolites due to the catalytic properties in fluid catalytic cracking, and organic synthesis and separation. Zeolite Beta always exists as an intergrowth of two end-member structures of polymorph A (*BEA, P4122 or P4322) and polymorph B (C2/c). Stacking faults dominate all over the Beta crystals and even some large pore defects were observed[1-2]. Due to this reason, no direct structure determination could be performed until recently. We reported a complete structure determination of polymorph B by the electron crystallography method on a polymorph B enriched zeolite Beta sample[3]. The polymorph B enriched zeolite Beta crystals show a wedge-shaped rod-like morphology. The surfaces of the crystals are not smooth but ridge-like as shown in Figure 1. HRTEM images were taken along the [1-10] direction to study the stacking behaviors in this material. A TEM image in Figure 2 shows that the stacking of 12-ring channels mostly follows the ABC or CBA type stacking mode of polymorph B inside the crystal (see Figure 3). These two types of stacking can be considered as two different twin components of polymorph B. The ABAB stacking of polymorph A (*BEA) is only observed at the twin boundaries which always run through the ab plane. It is also clear that there are no large pore defects inside the crystal; instead they are found near the surface of the crystal (Figure 4). Two large pore defects, one between the centre and the right ridge, and one between the center and the left ridge can be observed in Figure 4. It is very obvious that the ridge at centre shows a different stacking mode compared with the other areas, as indicated as white lines in Figure 4. The CBA type stacking in the center ridge and the ABC type stacking for the other areas belong to different polymorph B twin components, all of them grow from the same crystal ab plane. Since their frameworks do not fit each other, so they have to produce large pore defects in between. This significantly increases the lattice energy and may speeds up the dissolving. Due to the symmetry of the polymorph B structure, there should be also stacking faults in the orthogonal [110] direction. The stacking behavior at the crystal
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surface should be also the same. As we can find from Figure 1, the ridges can be observed along both two directions. 1. 2. 3.
J.M. Newsam, M.M.J. Treacy, W.T. Koetsier, C.B. de Gruyter, Proc. R. Soc. Lond. A. 1988, 420,375-405. Paul A. Wright, Wuzong Zhou, Joaquin Pérez- Pariente and Mar Arranz J. Am. Chem. Soc. 2005, 127,494-495. Corma, A., Moliner, M., Cantín, A., Díaz-Cabañas, M.J., Jordá, J.L., Zhang, D.L., Sun, J.L., Jansson, K., Hovmöller, S., Zou, X.D. Chem. Mater. 2008, in press.
Figure 1. A SEM image shows the morphology of polymorph B enriched zeolite Beta crystals. Insets are TEM images showing that the crystals are wedge-shaped in the [1-1 0] projection (top-right) and rectangular in the ab-plane (bottom-left). From Ref 3.
Figure 2. A HRTEM image taken along the [1-10] direction shows the ABC and CBA type stackings. Insets are Fourier transforms of (a) the whole image and (b) the pure ABC type stacking area. The ABC and BCA stacking areas are marked by white and black line, respectively.
Figure 3. Electron diffraction pattern and structure model viewed along the [1-10] direction of (above) a perfect polymorph B and (below) a twinned polymorph B.
Figure 4. A HRTEM image enlarged from the bottom left area of Figure 2 showing three ridges with different stacking modes. Large pore defects are observed between these ridges.
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Optimizing electron diffraction and EDS for phase identification in complex structures: application to multilayered Ti-Ni-P coatings P.A. Buffat1 and A. Czyrska-Filemonowicz2 1. Centre Interdisciplinaire de Microscopie Electronique, Ecole Polytechnique Fédérale de Lausanne, station 12, CH-1015 Lausanne Switzerland 2. Faculty of Metals Engineering and Industrial Computer Science, AGH University of Science and Technology, Al. Mickiewicza 30, 30-059 Kraków, Poland [email protected] Keywords: Ti-Ni-P, coatings, phase identification, diffraction, EDS, channeling
Electron diffraction, high resolution transmission electron microscopy (HRTEM), X-ray dispersive (EDS) and electron energy loss (EELS) spectroscopies constitute a set of well established techniques for sub-micrometer phase identification. All can now be incorporated in a single "analytical transmission electron microscope" and applied to the same area of investigation. However each of these techniques still require proper conditions to provide accurate information and using the same microscope settings, sample orientation and thickness for all of them lead most often to biased results. Electron diffraction retains the information of each phase or grain even when the electron beam reaches several of them at a time. Those lying close to a zone axis give rise to 2-D periodic lattices that can be easily distinguished from other reflections. Coupled effects between them add double diffraction reflections that are most often easily identified as such and coherent scattering by the overall entities is weak enough to be neglected if even noticeable. In that sense, electron diffraction gives unambiguous information on crystallographic data, but phase identification still requires referencing to databases where crystal symmetry and lattice parameters are linked to chemical composition. In the general case, the lattice spacing accuracy or full space group knowledge required to remove ambiguities between compounds of similar structure may be too high for electron diffraction techniques or too time consuming to gather, and eventually several possible solutions remains that can only be distinguished by an a priori knowledge of the chemical elements present or microanalysis (EDS, EELS). In contrast, EDS and EELS spectra from several phases lying simultaneously under the electron beam cannot be deconvoluted. The only chance lies in reducing the electron beam (probe) diameter with the hope that phases are not superimposed along the electron path. Then a qualitative analysis (which elements are present) seems straightforward. However when analyzing a minor phase in presence of a major one, the scattered electrons at the (thin) analyzed location hit (thick) parts of the sample far away after backscattering at pole-pieces or spiralling in the objective lens magnetic field (Figure 1). Their contribution adds a bias mainly representative of the major phase to the composition of the minor one that can amount up to some 10-20 at% [1]. This effect worsens with the sample tilt (tilting toward the EDS detector, as recommended to avoid shadow effects from sample holder in side entry goniometers, should be limited as much S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 681–682, DOI: 10.1007/978-3-540-85226-1_341, © Springer-Verlag Berlin Heidelberg 2008
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as possible!), but vanishes with FIB lift-off or omniprobe samples where the thick parts are removed. Figure 2 compares EDS analysis of a phase in a Ti-Ni-P coating on Ti6Al4V [2] from a classical (ion milled) cross-section to that one from a FIB lamella. Another important source of chemical analysis inaccuracy comes from the use of the same sample orientation for diffraction or HRTEM and EDS or EELS. Electrons travelling close to a zone axis or a strong diffraction direction suffer from important channeling effects that enhance or reduce the electron intensity, i.e. the ionisation efficiency, on specific atom columns (Figure 3). This effect may bias apparent compositions by several tens % or even more in particular cases. It is reduced by the use of a strongly convergent electron beam (large second condenser aperture), but requires tilting the sample quite far from the zone axis to become negligible [3]. 1. 2. 3. 4.
P.A. Buffat, EMS Yearbook (2004), p.78-90, http://infoscience.epfl.ch/record/102022/files/ A. Czyrska-Filemonowicz and P.A. Buffat, Micron (2008) in press K. Leifer et al, Micron 31 (2000), p. 411; J.C.H. Spence et al, Ultramicrosc. 26 (1988), p.103 Authors kindly acknowledge Dr. H.-J. Penkalla (Forschungszentrum Jülich) for supplying the FIB samples
Figure 1. EDS analysis bias due to scattered electrons toward thick sample parts.
Figure 2. EDS composition of the L5 sublayer in a Ti-Ni-P coating on Ti6Al4V [2] (a). Ti overestimation on classical cross-section leads to Ni(Ti0.93Al0.07)3.1 (b). Ni(Ti0.88Al0.12)2.1 from FIB lamella agrees well with NiTi2 found by SAED diffraction but Cu and Ga appear from the omniprobe support and ion implantation (c).
Figure 3. Channeling in Al.45Ga.55As. Integrated electron wave intensity along the electron path for zone axis [1,1,0] (a) and close to [120,119,-6] corresponding to kinematic (2,2,0) (b). Corresponding distributions of integrated intensity around atom columns at 100nm depth (red circles Al+Ga, yellow circles As) (c,d).
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Advanced analytical transmission electron microscopy to investigate the nano-graded materials properties M. Cheynet1, S. Pokrant2, L. Joly-Pottuz3, J.M. Martin4 1. SIMaP 1130 rue de la Piscine BP75 Saint Martin d’Hères 38042 France - 2 Carl Zeiss NTS GmbH, 73447 Oberkochen, Deutschland - 3. MATEIS, INSA, La Doua 69621 Villeurbanne, France - 4. LTDS, Ecole Centrale de Lyon, 69134 Ecully, France [email protected] Keywords: graded materials, EELS, structural, chemical, electronic and mechanical investigations
The concept of Graded Materials (GM) was proposed during the eighties. Since this period, a lot of experimental and theoretical research has demonstrated that GM offer properties that cannot be expected using homogeneous materials. In coated or composite materials, interfaces between different media are marked by sharp transition. These interfaces are thus the areas of high local stresses and adhesion problems. In GM, on the contrary, there is no singularity point (i.e. interface), but a continuous or in discrete steps change (in microstructure, composition, structure) within the volume. In the early time, research on GM has been widely focussed to optimize mechanical and thermal properties in structural materials [1]. Nowadays, with currently available new synthesis and processing capabilities (vapour or atomic layer deposition, molecular beam epitaxy, implantation, ion exchange, laser surface treatment), materials can be obtained with gradients at a nanometer length scale. They are consequently become more and more attractive for functional applications in disciplines as diverse as microelectronic, optoelectronic, high density storage media, tribology, catalysis or biomechanics. In consequence, high spatially resolved methods are required to investigate these NanoGraded Materials (NGM) and measure the impact of the gradation on their properties. Analytical Transmission Electron Microscopy is, of course, one of the most relevant [2], since it allows nowadays to perform imaging (HREM, HAADF, EFTEM) and spectroscopy (EDXS, EELS) with spatial and energy resolution at sub-nm and sub-eV level respectively. It is nowadays clearly demonstrated that TEM imaging combined with EELS analysis allow the mapping of electronic, dielectric, optical, chemical and mechanical properties at a nanometer scale [3]. This presentation will illustrate the relevance of Valence Electron Energy Loss Spectroscopy (VEELS) coupled to structural and chemical imaging to investigate and tailor functionally NGM properties. 1 - Structure gradation dependence of the dielectric properties of ALD HfO2 thin films [4]: Hafnium dioxide is one the most promising candidate to replace SiO2 gate oxide in the devices of the future technology nodes. Using HfO2, a 4 nm gate thickness would be adequate to prevent from current leakage. Therefore, in layer as thin as 4 nm, according to the deposition conditions, gradations in crystal structure occur and can significantly modify the dielectric properties. The best dielectric properties are obtained for HfO2 layer nano-crystallized in an orthorhombic structure.
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Sample
Structure
Eg(eV)
Si/SiO2/HfO2/Si Si/SiO2/HfO2/Ge
Ortho Mono
5.8 5.1
ε 5 4
Figure 1. Crystal structure of an ALD HfO2 layer capped in poly-Ge.( HREM images JEOL 400kV:J.L Rouvière CEA Grenoble). Si/SiO2/HfO2/poly-Ge 80% of monoclinic - 20% orthorhombic Table I. Band gap and electronic dielectric constant deduced from VEELS analysis. (VVELS experiments: TECNAI F20ST F. Tichelaar Delft).
2 - Carbon content gradation effect on the tribological properties of lubricant [5]. Diamond like carbon films are presently extensively studied for their tribological properties. ta-C (tetrahedral carbon) containing more than 70% of sp3 hybridized carbon is of particular interest because of its mechanical properties (hardness, Young’s modulus). When lubricated by alcohols, these coatings lead to a drastic reduction of friction coefficient never observed before. These outstanding properties were first attributed to a hydroxylation of their extreme surface. EFTEM and VEELS performed on FIB samples of ta-C coatings before and after friction show that these properties are also due to a change in the mechanical properties of the coatings. Since the coatings are amorphous materials without defects, it is possible to evaluate their hardness from the empiric relation proposed by Oleshko which correlates the hardness to the plasmon peak maximum energy [6]: log (H) =-7.44+6.1 log (Ep). Before friction, the coating is homogeneous (hardness of 52 GPa in the whole coating). After friction, a gradient of hardness is observed from 26 GPa at the surface to 43.7 GPa near the substrate. The changes of hardness can be explained by an accomodation of the coating during the friction and a transformation of sp3 in sp2 carbon. Figure 2. EFTEM image at 30 eV of the coating after friction test (left) VEELS spectra of a set of images between 13 and 40 eV (image-spectrum method). Ep is found to decrease from the substrat to the extreme surface indicating a change in the properties (hardness, Young’s modulus) of the coating during friction (right). 1. 2. 3. 4. 5. 6.
A. Mortensen and S. Suresh, Int. Mater. Rev. 40, 239, (1995). D. A. Muller, T. Sorsch, S. Moccio, F. H. Baumann, K. Evans-Lutterodt, G. Gimp, Nature, 399, 758, (1999). R. Erni, N. Browning, Ultramicroscopy, 107, 267, (2007). M. C. Cheynet, S. Pokrant, F. Tichelaar, J. L. Rouvière, J. Appl. Phys. 101, 054101 (2007). L. Joly-Pottuz et al., Journal of Applied Physics 102, 064912 (2007). J. M. Howe and V. P. Oleshko, Journal of Electron Microscopy 53, 339 (2004).
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Characterisation of Nickel Nanocomposites by SEM, TEM and EBSD D. Dietrich, Th. Lampke, B. Wielage, D. Thiemig1 and A. Bund1 Institute of Composite Materials and Surface Technology, Chemnitz University of Technology, D-09107 Chemnitz, Germany 1. Department of Physical Chemistry, Dresden University of Technology, D-01062 Dresden, Germany [email protected] Keywords: Electroplating, nanoparticles, microstructure.
Electroplated composites with incorporated nanoparticles become a key material in the fabrication of microsystem devices. A good survey of composites with nickel matrix is given in [1]. Electrodeposition of alumina particles (13 nm diameter, content around 1 Wt %) in nickel from an acidic electrolyte has been investigated using a nickel sulfamate bath with a particle loading of 10 g/l. Direct current (DC) deposition was compared to pulse-reverse plating (PRP) which can facilitate higher amounts. Details are given in [2]. Materials properties depend on the particle distribution, their bonding to the matrix and their influence on the microstructure of the matrix which has been characterized by TEM, STEM, SEM and EBSD. A FESEM NEON40EsB was used at 25 kV with a STEM detector as well as an EBSD camera (EDAX TSL). For QBSD and SE imaging, the voltage was lowered to 10 and 5 kV, respectively. Cross sections were prepared with a final OP polish. For TEM (HITACHI 8100, LaB6 cathode, 200 kV) studies, 3 mm disks were grinded, cut and polished by 3 kV Ar ions with 6° incidence angle. Starting with a nearly 1 µm thick fine crystalline initial layer the subsequent layer is formed by columnar crystals with growing size from 0.4 µm to 1.2 µm and a <100> fibre texture in growth direction. This is reflected in the orientation contrast images (Figures 1-3). The alumina particles agglomerate (Figure 4) and form chains which can easily be detected in STEM dark field imaging (Figure 5). Pulse-reverse plating seems to promote the columnar growth, twin formation and a sharper texture (Figure 6 a-b). In the IPF colour coded nickel map (Figure 6c) the black spots represent highly disordered matrix parts containing particles. Compared to STEM imaging particles can hardly be distinguished from the matrix in TEM (Figure 7). Nevertheless, a good particle matrix bonding without voids can be confirmed only by TEM at higher magnification (Figure 8). Such studies reveal that agglomerated particles are preferentially incorporated in grain and twin boundaries. 1. 2.
C.T.J. Low, R.G.A.Wills, F.C.Walsh, Surf. Coat. Technol. 201 (2006) 371-383. D. Thiemig, R. Lange, A. Bund, Electrochim. Acta 52 (2007) 7362-7371.
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Figure 1. DC nickel alumina composite.
Figure 2. PRP nickel layer.
Figure 3. PRP nickel alumina composite.
Figure 4. Particle agglomerates.
Figure 5. Particle distribution in STEM.
Figure 6c. Ni IPF map
Figure 6 a-b. Pole figures of DC plated (a) and PRP (b) composites.
Figure 7. Particle chains.
Figure 8. Detail.
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Characterisation of Gold Nanocomposites by SEM, TEM and EBSD D. Dietrich, Th. Lampke, B. Wielage, P. Cojocaru1 and P.L. Cavallotti1 Professur Verbundwerkstoffe, Technische Universität Chemnitz, D-09107 Chemnitz, Germany 1. Dipartimento di Chimica, Materiali, Ingegneria Chimica, Politecnico di Milano, 20131 Milano, Italy [email protected] Keywords: Electroplating, nanoparticles, microstructure.
The electrodeposition of gold is a key technology in the fabrication of many microsystem devices comprising gold composites with improved electronic and mechanical properties [1]. Multiwalled carbon nanotubes (CNT) of 20-40 nm in diameter were obtained by catalytic decomposition of acetylene on an iron catalyst. Ultradispersed diamond (UDD) was obtained by detonation synthesis (Caspio SA, Switzerland). Gold composites were electroplated using a sulphite gold electroplating bath with a CNT loading of 3 g/l and a UDD loading of 20 g/l, respectively [2]. The particle distribution, the bonding to the matrix and the influence on the matrix microstructure were characterized by complementing electron-microscopic methods. A FESEM NEON40EsB was used at 25 kV with a STEM detector as well as an EBSD camera (EDAX TSL). For EsB and Inlens SE imaging, the voltage was lowered to 1 kV. Cross sections were prepared by diamond grinding with a final Ar ion (2.5 kV 15°) polish. For TEM (HITACHI 8100, LaB6 cathode, 200 kV) studies, 3 mm disks were cut, grinded and thinned to electron transparency by Ar ions (3kV, 6°). Pure gold with a grain size of around 2 μm reveals characteristically large twins as well as series of twin lamellae down to 5 nm in width. The Au/CNT composite shows a reduced grain size of 400 nm with a strikingly pronounced twin formation. This is associated with CNT bath loading; however, CNTs could not be confirmed either by TEM or by EsB. The UDD gold composite reveals a grain refinement down to 50 nm (Figures 1a-c and 2a-c). Mean grain sizes and film textures were derived from EBSD. All films show a {111} fibre texture in film growth direction (Figure 3). Numerous incorporated diamond nanoparticles with a diameter down to 5 nm were detected by SE (Figure 4), EsB and STEM (Figure 6) which show an even distribution without agglomeration. By HRTEM, the gold matrix with {111} lattice planes (0.235 nm spacing) and {200} lattice planes (0.204 nm spacing) can be distinguished around the nanodiamonds, which shows a proper incorporation. Nevertheless, no particle lattice order could be observed. 1. 2.
H.Y. Song, X.W. Zha, Physica B 403 (2008) 559–563. P. Cojocaru, A. Vicenzo, P. L. Cavallotti, J. Solid State Electrochem. 9 (2005) 850-859.
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Figure 1a-c. BS orientation contrast: Pure Au (a), Au/CNT (b) and Au/UDD (c).
Figure 2a-c. IPF orientation map: Pure Au (a), Au/CNT (b) and Au/UDD (c).
Figure 3. Pole figure of AU/CNT.
Figure 4. InLens SE: Au/UDD.
Figure 5. Diamond particle in gold matrix.
Figure 6. STEM: UDD distribution.
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Alumina Coatings as Protection against Corrosive Atmosphere I. Dörfel, R. Sojref, M. Dressler, D. Hünert, M. Nofz Federal Institute for Materials Research and Testing (BAM) Unter den Eichen 87, 12205 Berlin, Germany [email protected] Keywords: alumina, protection layers, microstructure
Oxidation protection layers were deposited by spin coating on substrates of Ni based superalloy SC 16 [1]. The protection effect of that type of coatings during heat treatment in dry air at 800°C for 4000 hours was investigated by Dressler [2]. In order to develop protection coatings also effective in a steam containing corrosive atmosphere, different types of multilayered coating systems were tested in an environment of a flowing mixture of CO2 and steam at 500 °C for 120 hours and subsequently at 700 °C for another 120 hours. Here the results of the investigations at two different types of coating systems will be presented: (i) two layers of a mixture of ethanolic boehmite sol with a suspension of submicron corundum powder and (ii) the two layers of (i) covered by a third layer of an aqueous sol. Before testing in corrosive atmosphere both types of samples were heat treated at 500 °C for 0.5 hours and at 800 °C for 2 hours in air. The cross sectional TEM samples were prepared by the FIB technique in a Strata TM 200 xP (FEI) and investigations were carried out in an analytical STEM JEM 2200FS (JEOL) equipped with a FEG and an omega filter. TEM and STEM methods were combined with electron diffraction investigations and EDX spectroscopy in order to characterize microstructure and composition of the tested samples. The influence of the corrosive atmosphere is significant. The sol components of the sol/suspension layers are expected to form a fine crystalline fraction around the corundum grains, comparable with that of the same coating type heat treated in dry air at 800 °C for 4000 hours (Figure 1). During the steam testes this fine crystalline fraction of the microstructure is reduced by grain coarsening. Large grains were embedded in residues of that fine crystalline matrix containing different holes. The large grains were not only corundum added to the sol but a part of them was identified as δ- Al2O3 grains like the fine crystalline fraction too. During their crystal growth, the larger grains consumed the fine crystalline matrix, gaps occurred and in the course of the sample preparation loose grains turned off. The effect of grain coarsening is more pronounced in the sample (i). The seeding effect of the corundum grains and grain growth are accelerated by action of the steam/CO2 mixture. This is in agreement with [3]. The cover layer of mere sol on sample (ii) has temporary protection effects against the steam. In the region of the cover layer no large grains but a fine crystalline microstructure originated, consisting of nano-sized particles of the δ- Al2O3 phase too. Beside the coarsening effects in the sol/suspension zones diffusion processes mainly of
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the elements Cr, Ti and Ni from the substrate into the coating are observed. Because of the relative moderate temperatures and short times during the steam exposure the diffusion zone is limited to the region near the interface substrate/coating. 1. 2. 3. 4. 5.
M. Nofz, I. Dörfel, R. Sojref, Thin Solid Films 515 (2007) p.7145 M. Dressler, PhD Thesis, TU Bergakademie Freiberg, Germany 2007 R. B. Bagwell, G. L. Messing J. Am. Ceram. Soc. 82 (1999) p. 825 The authors wish to thank H. Rooch and W. Gesatzke for preparing the TEM lamellae. The project was financially supported by the Deutsche Forschungsgemeinschaft.
Figure 1. STEM BF image of a sol/suspension alumina coating heat treated at 800 °C for 4000 hours
Figure 2. TEM micrographs of the samples i) (left) and ii) (right hand side).The large grains and some holes are clearly visible, especially in i), whereas the cover layer of ii) is nano crystalline and non-degenerated and the large grains have lower dimensions
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Advanced Multilayer Systems for X-ray Optics: Quality Assessment by TEM D. Häussler1, W. Jäger1, E. Spiecker1, B. Ögüt1, U. Ross1, J. Wiesmann2, M. Störmer3 1. Microanalysis of Materials, Faculty of Engineering, Christian-AlbrechtsUniversity of Kiel, 24143 Kiel, Germany 2. Incoatec GmbH, 21502 Geesthacht, Germany 3. Institute of Materials Research, GKSS Forschungszentrum Geesthacht GmbH, 21505 Geesthacht, Germany [email protected] Keywords: multilayers, X-ray optics, TEM, Geometric Phase Analysis
Precisely engineered systems of nanoscale multilayers are essential device components in X-ray optics for spectrometry and in synchrotron applications. Their fabrication and optimization require processing control by high-resolution microstructure characterization. Methods of TEM have proven to be indispensable in quantitatively assessing properties like multilayer periodicity, orientation, or interface roughness, and in correlating the microstructure with X-ray reflectivity [1, 2]. Examples for such characterisations are given for advanced multilayer systems on Si gratings and for multilayer system with depth gradients. Such systems are used as spectral filters for low-bandwidth and high-bandwidth applications, respectively. A bright-field (BF) crosssection image analysis of a Mo-B4C multilayer system on a Si grating (Figure 1) reveals regions of high multilayer perfection and regions of imperfect growth. The application of the geometric phase analysis (GPA) [1] to cross-section TEM micrographs of a W-C multilayer system with depth gradient is shown in Figure 2. This method enables the quantitative assessment of the local multilayer perfection and overall homogeneity. Further examples are extended multilayer systems with lengths > 1000 mm for large X-ray optics used in synchrotron applications. 1. 2.
D. Häußler, E. Spiecker, E. Yang, W. Jäger, M. Störmer, R. Bormann, G. Zwicker, phys. stat. sol. 202 (2005), p. 2299. D. Häußler, E. Spiecker, W. Jäger, M. Störmer, R. Bormann, C. Michaelsen, J. Wiesmann, G. Zwicker, R. Benbalagh, J.-M. André, P. Jonnard, Microelectronic Engineering 84 (2007) p. 454.
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Figure 1. Mo-B4C multilayer system on a Si grating. Cross-section BF-TEM micrograph (left). Magnified section (right) shows perfect multilayer and edge region.
Figure 2. W-C multilayer system with depth gradient: stacks of 15 bilayers, bilayer thickness range 3nm -6nm. Cross-section BF-TEM (left), local homogeneity of bilayer thickness as revealed by geometric phase analysis (false-color representation).
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Surface investigation of SU-8 by atomic force and scanning electron microscopy Th. Kups1, Chr. Kremin2, M. Hoffmann2 and L. Spieß1 1. Institute of Micro- and Nanotechnologies, Materials of Electronics, Ilmenau University of Technology, P.O. Box 100565, 98684 Ilmenau, Germany 2. Institute of Micro- and Nanotechnologies, Micromechanical Systems, Ilmenau University of Technology, P.O. Box 100565, 98684 Ilmenau, Germany [email protected] Keywords: SU-8, AFM, SEM
SU-8 is a recently developed negative, epoxy-type, near-UV photoresist. It offers good mechanical strength and chemical resistance. Once it is cured, it is essentially impossible to remove without aggressive applications. Moreover, this material is sufficiently cheap to be considered in low cost (or even disposable) devices. Therefore, SU-8 photoresist has a wide range of possible applications: microelectronics, micromechanics, microfluidics, packaging, magnetics and others [1, 2, 3]. Additionally an increasing interest in using the SU-8 for producing bio-MEMS or microfluidic devices is noticeable. In biological applications, sterilisation processes using chemical treatments or UV radiation are common practice and the analysis of their interaction with SU-8 is of great interest [4]. In this work AFM and SEM investigations will be presented to show the changes in morphology after strong / long UV exposition in addition to several post processing steps of SU-8 layer formation on glass substrate. Layers of photoresist SU-8 (GM1070 by Gersteltec GmbH; SU-8 2075 by Microchem Corp.) were made by spin-coating a 50-100 µm thick resist layers onto glass substrates at 1000 rpm for 30 s. After a soft-bake step the coatings were finally exposed by UV-light of two different doses of 505 mJ/cm² for 210 s and 720 mJ/cm² for 300 s. For stability investigations the specimens were on the one hand hard-baked (150°C for 30 min) or strongly UV exposited (1440 mJ/cm² for 10 min) and on the other hand stored for 7 days at standard room conditions to test the long time stability of the layers. The structures were then analysed by scanning electron microscopy (SEM) and atomic force microscopy (AFM). For SEM investigation the SU8 layers were additionally covered by a 15-20 nm carbon film which was sputtered onto the surface to increase electrical conductivity. A FEI XL30 SEM operating at 5-10 kV (to avoid charging effects) in SE – mode was used for investigation. AFM measurements were done in noncontact mode by using an Atos Solver Pro (NT-MDT) AFM. In Figures 1 and 2 shows SEM and AFM images of SU-8 layers after coating (Fig. 1) and after storage of several days (Fig. 2). After coating and short post-bake process the surface shows many cyclic structures (Fig. 1) which has a RMS of about 50 nm what can be interpreted as leftover of solvents by the coating process. SU-8 of this sample has a pitted and porous surface structure. Contrary to that after long time storage for
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seven days at room temperature and pressure and at day light the surface shows a complete different structure (Fig. 2). The surface shows not longer cyclic structures but cracks with depth of about 20-30 nm in all magnification ranges of investigation. The low dimensional surface structure (1 µm x 1 µm) shows instead of a pitted and porous structure a smoothed and more uniform surface structure.
Figure 1. SEM and AFM images of a sample after coating with UV exposition of 50 µm at 505 mJ/cm² for 210 s show cyclic structures and porous surface
Figure 2. SEM and AFM images of a sample for long time stability test (7 days) at standard room conditions show strong breaks and changes of the surface structure
1. N.-T. Nguyen and S.T. Wereley, Fundamentals and Applications of Microfluidics, Artech House, 2002. 2. B.A. Auld, Acoustic Fields and Waves in Solids, vol. II, Krieger, Florida, 1990 3. N. Arana, D. Puente, I. Ayerdi, E. Castano and J. Berganzo, Sensors and Actuators B 118 (2006), p. 374–379 4. A. Keppler, M. Himmerlich, Chr. Kremin, J. T. Schumacher, A. Grodrian, J. A. Schäfer, J. Metze, M. Hoffmann and St. Krischok, Frühjahrstagung der DPG Berlin 2008 5. This project was kindly supported by the Thuringian Ministery of Culture within the project “Intergrierte Mikrooptische Pinzette”
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SEM and TEM investigations of electrophoretical deposited Si3N4 and SiC particles in siloxane of steel substrate Th. Kups1, A. Knote2 and L. Spieß1 1. Institute of Micro- and Nanotechnologies, Materials of Electronics, Ilmenau University of Technology, P.O. Box 100565, 98684 Ilmenau, Germany 2. Department of Metallic Materials and Composite Materials, Ilmenau University of Technology, P.O. Box 100565, 98684 Ilmenau, Germany [email protected] Keywords: electron microscopy, electrophorese, particles
Metal-ceramic composite coatings on metal structural components are of great technical interest, primarily for wear and corrosion protection. These coatings are expected to exhibit both a high wear and corrosion resistance, a good adhesive bonding to the metallic phase as well a high damage tolerance during mechanical stresses. A possibility to produce uniform metal-ceramic composite coatings with a high content of ceramic particles will be presented in this study [1, 2, 3]. This method includes a combination of electrophoretic deposition and electrolytic deposition by several steps. A ceramic coating (SiC or Si3N4) was electrophoretically deposited from a aqueous suspension on a ferritic steel plate. For the suspension firstly was attenuated commercial siloxane (Dynasilan HS 2909; Degussa) by de-ionized water down to a concentration of 3% of the active agent. In this aqueous solution the ceramic particles (content of the solid: 12%) are dispersed ultrasonically to an agglomerate-free suspension. The electrophoretic deposition was realised at current densities of 0.2 – 2.0 mA/cm² at coating times between 3 and 10 s. After drying the layer in air, this was sintered at temperatures between 600°C and 800°C for 1 h to an open porous layer. By changing the sintering temperature the porous structures can be changed. The resulting layers have a thickness between 10 – 30 µm. The structural characterisation of the ceramic layers was performed by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). For the SEM investigation particles from the top layer region were mechanically removed and inner layer particles were put onto a SEM holder. To prepare a TEM sample particles were mechanically removed from the inner layer regions crushed by a motar to break the clustering of the particles and put on a gold grid covered with a about 10 nm thick carbon film. Fig. 1 shows SEM images of open porous ceramic layers with SiC particles partly covered by siloxane to bond one particle with another particle a) specimen sintered at 600°C and b) sintered at 750°C. It can be clearly seen for the case of the low temperature sintering the bonding between the siloxane and the particle is stronger and the siloxane bridges are bigger than in the high temperature case where only small and flat bondings between the particles are realised by the siloxane. Fig. 2 shows the high resolution TEM images of the interface between the Si3N4 particle and the siloxane. Because S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 695–696, DOI: 10.1007/978-3-540-85226-1_348, © Springer-Verlag Berlin Heidelberg 2008
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siloxane has an amorphous structure the particles show a non-crystalline border with a size of 5-10 nm.
Figure 1. Scanning electron microscopy images of SiC particles partly covered by siloxane sintered at a) 600°C and b) 750°C. In a) the bonding between the siloxane and the particle is stronger and the siloxane bridges are bigger than in case b) where only small and flat bondings between the particles are realised by the siloxane.
Figure 2. High resolution transmission electron microscopy images show the interface between the Si3N4 particle and the siloxane sintered at 800°C. The siloxane has an amorphous structure therefore the particles show a non-crystalline edge with a size of about 5 nm around the crystalline particles. 1. W. Burghardt, Galvanotechnik 85 (1994), 406 2. C. Karayianni, and P. Vassiliou, J. Mater. Sci. Lett 17 (1998), 398 3. A. Knote, H. G. Krüger, S. Selve, Th. Kups, H. Kern and L. Spiess, J. Mater. Sci 42 (2007), 4545
4. This project was kindly supported by the German Research Foundation by the project KE 359/8-1
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Contribution of electron microscopy techniques to the chemical and structural characterization of TiC/a-C nanocomposite coatings C. López-Cartes, D. Martínez-Martínez, J.C. Sánchez-López and A. Fernández Instituto de Ciencia de Materiales de Sevilla (Univ. Sevilla-CSIC), Avda. Américo Vespucio nº49, 41092 Seville (Spain) [email protected] Keywords: Nanocomposite, coating materials, TiC, amorphous carbon, hardness, friction
The design of coatings with nanocomposite structure, consisting of at least two phases (one or two that are nanocrystalline and one amorphous), is attracting an increasing interest due to the new possibilities of synthesis of applicable materials with superior physical, chemical and engineering properties [1]. In the field of the design of hard coatings and their application to protect cutting or drilling tools, the search for new systems with improved mechanical and tribological properties is highly demanded. The knowledge of the structure of the nanocomposites is considered of great importance not only for understanding the mechanical behaviour at the microscale, but also for the design of strategies to synthesize materials with controlled properties. In this work the results concerning the synthesis and characterization of a series of TiC/a-C nanocomposite coatings prepared by magnetron sputtering of titanium and graphite targets are presented. The coatings have been chemical and structurally characterized by using a battery of electron microscopy techniques (TEM, ED, HRTEM, SEM and EELS) in order to establish correlations between the synthesis conditions, their structure and mechanical and tribological properties. The structure and the compositions of the coatings are very dependent of the synthesis conditions, and basically related to the ratio of power applied to each target (WC/WTi) during the preparation. As revealed by HRTEM studies, the nanostructure of the coatings changes from quasipolycrystalline TiC, consisting of small crystals (7-10 nm) surrounded by a thin layer of amorphous material in the case of the samples prepared using WC/WTi=1, to nanocomposites formed by nanocrystals of TiC (1-2 nm) embedded in an amorphous matrix of carbon for WC/WTi=4 (Figure 1). Using linear combinations of EELS spectra recorded at K-edge of carbon from two reference samples (pure TiC and amorphous carbon), an enrichment in amorphous carbon and a diminution of TiC crystalline phase from 95 to 25% has been found in parallel to the increment of WC/WTi from 1 to 4 [2]. Cross section images obtained by SEM clearly demonstrate that the microstructure of the coatings is also affected by the preparation conditions. Columnar structure is observed when low power ratios are used and a more compact structure is achieved when coatings are synthesized at higher power ratios. Both mechanical and tribological properties of the coatings are dependent on the structure and chemical composition. The coatings with high content in crystalline TiC show the better hardness values (27 GPa) but poor tribological behaviour. As the
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content in amorphous carbon is increased, the friction coefficient is diminishing to values equal to 0.1 approximately and the coatings present lower wear rates (k~mm3/Nm). However, the softer nature of a-C leads to lower hardness values (5-10 GPa). 1. 2. 3.
A.A. Voevodin, S.V. Prasad and J.S. Zabinski, J. Appl. Phys. 82 (1997) 855. D. Martínez-Martínez, C. López-Cartes, A. Justo, A. Fernáncez, J.C. Sánchez-López, A. García-Luis, M. Brizuela and J.I. Oñate, J. Vac. Sci. Technol. A, 23(6) (2005) 1732. We acknowledge the financial support from the Spanish MEC (project no. MAT2004-01052 and MAT2007-66881-C02-01).
Figure 1. High-resolution images of coatings prepared using power ratios of WC/WTi=1 (a) and WC/WTi=4 (b). The digital diffraction pattern of the central crystal in (a) is inserted in the image.
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TEM investigations of the Ti/TiN multilayered coatings deposited on the Ti-6Al-7Nb alloy T. Moskalewicz1, H.J. Penkalla2 and A. Czyrska-Filemonowicz1 1. AGH University of Science and Technology, Faculty of Metals Engineering and Industrial Computer Science, 30-059 Kraków, Al. Mickiewicza 30, Poland 2. Forschungszentrum Jülich GmbH, D-52428 Jülich, Germany [email protected] Keywords: coatings, multilayer, titanium alloy, TEM, pulsed laser deposition (PLD)
Titanium alloys are known to be among the best biocompatible metallic materials, however they exhibit a relatively poor wear resistance [1]. One of the possibilities for improvement is to process thin coatings on their surfaces, e.g. by pulsed laser deposition technique (PLD) [2]. The goal of the present investigations was to characterise a microstructure of the Ti/TiN multilayered coatings formed on the Ti-6Al-7Nb alloy by the PLD. The Ti-6Al-7Nb alloy, delivered as annealed at 750°C/2 h, was subsequently surface treated by PLD technique at room temperature. The Ti/TiN multilayers coatings composed of 32 and 64 layers were produced. Details concerning coating processing are given in Ref. [3]. Microstructure and phase composition of the Ti/TiN multilayers were characterised by analytical TEM (SAED, EDS, EFTEM) on cross-section thin foils prepared by FIB. Analytical TEM investigations were performed using JEM-2010 ARP and PHILIPS CM200. Microstructure of the coating composed of 64 layers is shown on Figure 1. The SAED patterns, EDS and EFTEM analyses were used for phase identification. The SAED patterns were interpreted with the JEMS software [4]. Phase identification revealed that both coatings were consisting of α-Ti (hexagonal closepacked; hcp) and δ-TiN (face-centred cubic NaCl type; fcc) layers. The coatings were composed of columnar crystallites. The total thickness of the both coatings was constant and equal to 1.1 µm. Thickness of particular layers in the multilayered coatings was evaluated from TEM dark-field images and EFTEM elemental maps. The thickness of δ-TiN and α-Ti layers decreased with increasing of component layers number. In the case of the 32 layer system, it was 32 and 34 nm for δ-TiN and α-Ti, while 15 and 20 nm for 64 layer system, respectively. EFTEM elemental maps showed the presence of nitrogen in every second layer (Fig. 2). A relationship between the micro/nanostructure of the investigated multilayered material and it mechanical and tribological properties was established [3]. The coatings exhibited much higher microhardness than that of the substrate. Result of scratch-test showed that the 64 layer coating has better adhesion to the substrate than 32 layer one. The main failure mechanism of the coatings was a wear process. 1.
D.M. Brunette, P. Tengvall, M. Textor, P. Thomsen, Titanium in medicine. Berlin: SpringerVerlag; 2001.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 699–700, DOI: 10.1007/978-3-540-85226-1_350, © Springer-Verlag Berlin Heidelberg 2008
700 2. 3. 4. 5.
L. Major, J. Morgiel, B. Major, J.M. Lackner, W. Waldhauser, R. Ebner, L. Nistor, G. Van Tendeloo, Surf. Coat. Technol. 200 (2006) 6190. T. Moskalewicz, R. Seddon, R. Major, J.M. Lackner, A. Czyrska-Filemonowicz, submitted to Surf. Coat. Technol. P. Stadelmann, (2004) JEMS Java Electron Microscopy Software, http://cimewww.epfl.ch/people/stadelmann/jemswebsite/ jems.html. This work was supported in part by the EU Network of Excellence project Knowledge-based Multicomponent Materials for Durable and Safe Performance (KMM-NoE) under the contract no. NMP3-CT-2004-502243. The surface treatment was carried out in the Laser Center Leoben (MCL), Austria. The valuable contribution of Prof. J.M. Lackner (MCL) and Dr R. Major (IMIM PAN) is kindly acknowledged.
Figure 1. Microstructure of the Ti/TiN multilayer coating (64 layers) on the Ti-6Al-7Nb alloy and corresponding SAED pattern taken from the area marked as ring and its identification, TEM BF, cross-section FIB thin foil
Figure 2. a) TEM BF image of the Ti/TiN multilayer coating consisting of 32 layers; b) Ti L2,3 jump ratio image; c) N K jump ratio image
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Microstructural examination of Al and Cr alloyed zinc coatings on low carbon steels D. Chaliampalias1, G. Vourlias1, E. Pavlidou1, K. Chrissafis1, G. Stergioudis1, S. Skolianos2 1. Physics Department, Aristotle University of Thessaloniki, Greece 2. Department of Mechanical Engineering, Aristotle University of Thessaloniki, Greece [email protected] Keywords: Coatings, SEM, Pack Cementation, Alloying elements
Zn, Al and Cr are considered as well-suited metals for corrosion protection coatings on steels. Their performance is related to their ability to form dense, adherent films on the surface of the steels with substantial lower corrosion rate than the ferrous materials [1]. Additionally as Zn is anodic to the steel, can be cathodically protect it [1]. Moreover, Cr forms several hard phases which enforce the hardness of the surface [2]. In this work the feasibility and structure of zinc coatings alloyed with Al and Cr is examined. The coating was developed by CVD pack cementation technique [3] and its characterization was accomplished using a JEOL 840A SEM (20 kV) equipped with an ISIS 2000 EDS analyzer. Common Zn coatings (fig. 1) are composed by Zn-Fe layers as a result of the diffusion of Fe from the substrate into the as formed coating. In this phase, Zn concentration increases near the surface [3] and decreases at an amount lower than that of Fe, near the interface. Alloying zinc coatings with Al (fig. 2), results to an additional layer on the surface which is mainly composed of Fe, Al and small amounts of Zn (inset images of fig. 2). No other Al amounts were found in the coating. This can be attributed to the fact that during the thermal treatment, Al diffused towards the surface and formed the denoted layer in fig. 2. EDS line scanning revealed that the thickness of this layer was 25μm and it was formed over a 50μm Zn-Fe layer (fig. 3). The formation of this layer is expected to enforce the corrosion protection beyond the level offered by the underlying thick Fe-Zn coating. Similar results were observed in the case of Cr alloyed zinc coatings. Cr was gathered on the coating surface forming an additional layer (fig. 4). This layer contains Zn, Cr and some small amounts of Fe and is expected to enhance the corrosion resistance and increase the hardness of the specimen [2]. The thickness of this layer is proportional to Cr concentration in the pack mixture and at elevated Cr concentrations (20%wt) it was measured up to 90 μm (fig. 5). Below this layer the same phases were observed as in the case of pure Zn coatings, which also contribute to the substrate protection. 1. 2. 3.
Marder A. R., Prog. Mat. Sci., 45 (2000), p.191. Pistofidis, N., Vourlias, G., Stergioudis, G. Corrosion Engineering Science and Technology 42 (1) 2007, p. 16. Vourlias G., Pistofidis N., Chaliampalias D., Patsalas P., Pavlidou E., Stergioudis G. and Tsipas D., Polychroniadis E.K., Surf. Coat. Tech., 200 (2006), p.6594.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 701–702, DOI: 10.1007/978-3-540-85226-1_351, © Springer-Verlag Berlin Heidelberg 2008
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Study of the structure and high temperature oxidation resistance of high alloyed tool steels E. Pavlidou, D. Chaliampalias, G. Vourlias, K. Chrissafis Physics Department, Aristotle University of Thessaloniki, 54124 Thessaloniki, Greece [email protected] Keywords: Tool steels, Oxidation, Electron Microscopy, TG Measurements
Tool steels are widely used in several mechanical devices that demand materials with special properties. High temperature oxidation resistance of tool steels is required on many applications such as casting and forging dies, blades for hot shearing and hot extrusion tools [1-2]. In this work two high grade tool steels, AISI D6 (C:2.05%w.t., Si:0,30%w.t., Mn:0,80%w.t., Cr:12,50%w.t., W:1,30%w.t, Fe: Bal.) and Cr-Mo-V (C:0,60%w.t., Si:0,35%w.t., Mn:0,80%w.t., Cr:4,50%w.t., V:0,20%w.t, Mo:0,50%w.t. Fe: Bal), were structurally examined before and after being exposed in high temperature oxidative environment. Structural examinations were performed using a JEOL 840A SEM (20 kV) and a 100kV JEOL 100CX TEM. Oxidation tests were accomplished by thermogavimetric measurements with a SETARAM SETSYS TG-DTA 16/18. D6 tool steel is characterized by large-sized dark grey islands (fig.1a area 1) which contain 65-70%wt. Fe and 25-30%w.t. Cr which were recorded in the ED TEM pattern (fig. 1b). Light colored areas in the matrix contain small amounts of tungsten (fig. 1a, area 2) and common matrix areas contain Fe and 7%w.t. Cr (fig.1, area 3). Cr–Mo–V steels contain dispersed spherical agglomerations which contain 45-50%w.t.Cr and 58%w.t.V (fig.1c, area 1), also recorded in the ED TEM pattern (fig. 1d), while matrix areas (fig. 1c, area 2) contain Fe with low Cr concentration. After non isothermal oxidation from ambient temperature to 1000 oC, a 60μm oxide layer was formed on D6 steel (fig. 2a) while on Cr-Mo-V steel was measured 40μm (fig. 2b). Cross section EDS linear analysis showed that the upper part of both layers contain Fe oxide while the inner areas contain chrome oxide. The formation of chrome oxide layer is mostly due to high temperature, which made Cr amounts to rise up to the surface and react with the environment. This layer, on D6 steel, has elevated thickness (35μm) comparing with that of Cr-Mo-V steel (20μm). Considering that D6 steel contain more Cr, it explains the difference between its thicknesses on the herein examined tool steels. Non isothermal curves showed that Cr-Mo-V tool steel begin to oxidize earlier than D6 (fig. 3a). This is mostly due to the thinner chrome oxide layer in Cr-Mo-V tool steel which is known to be protective in high temperature environments. Plane view SEM micrographs demonstrate the tool steel surface morphology after high temperature oxidation test. Fe oxide on the surface of D6 steels (fig.3 b,c) is very compact while on Cr-Mo-V steel is porous (fig.3 d,e). 1. 2. 3.
E. Gemelli, A. Gallerie, M. Caillet, Scripta Mater. 39 (1998), p. 1345. Z.k. Yu, J.T. Lu, Surf. Eng. 3 (1) (1987), p. 41. ASM International, “Forms of Corrosion”, ASM handbook of metals, Vol.13, p.97.
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A comparative study of NiCrBSi and Al coated steels with thermal spray process in different environments D. Chaliampalias1, G. Vourlias1, E. Pavlidou1, K. Chrissafis1, G. Stergioudis1, S. Skolianos2 1. Physics Department, Aristotle University of Thessaloniki, Greece 2. Department of Mechanical Engineering, Aristotle University of Thessaloniki, Greece [email protected] Keywords: Coatings,Oxidation, Thermal Spray, SEM, Thermogavimetric Measurements
The application of coatings is a well known method for steel protection because they act as a shield against different corrosive environments [1, 2]. The aim of this work is a comparative examination of the resistance of thermal sprayed NiCrBSi and Al coated low carbon steels in high temperature environments and in marine atmosphere. Microscopic observations were accomplished using a JEOL 840A SEM (20 kV). Oxidation tests were performed with a SETARAM SETSYS TG-DTA 16/18 apparatus and in marine simulated atmosphere in a salt spray chamber (SC-450). After non isothermal oxidation, from ambient temperature to 1000oC, chrome and silicon rose up to the surface from the inner areas of the NiCrBSi coatings forming a withstanding silicon and chrome oxide layer [3] (fig.1a, c). On Al coating, the surface contains only Al oxide and some Fe oxide compounds while Al-Fe phases were recorded both underneath the surface and into the substrate, in form of “paint runs”. This can be attributed to diffusion mechanisms activated by the high temperature (fig. 1b, d). TG curves (fig. 2a) verify that Al coatings oxidize earlier (650oC) than NiCrBSi (750oC). On the surface of NiCrBSi coatings several local cracks were found due to Cr evaporation [3] at high temperatures (fig. 2b) while the surface of Al coatings shows serious damages (fig. 2c). After 1 month of subjection in the salt spray chamber, two layers, corresponding to Fe oxide, were formed on the interface and on the top of the NiCrBSi coatings (fig.3a) and no Cl amounts were recorded. These layers were created by the diffusion of corrosive elements through the preexisting porous net in the flame sprayed coating. The formation of the Fe oxide on the interface, decreases the adhesion strength of the coating.The middle light colored areas, between the two Fe oxide layers, are mainly composed by Ni, Cr and Si. Al sprayed coatings contain small Cl amounts throughout the whole coating and are much more withstanding in such environments as they remain homogeneous with the exception of particular areas where O2 together with Al and Cl were recorded (fig.3b). The enhanced durability of Al coatings attributed to the formation of a bulk aluminum oxide thin film on the surface of the coating which acts as a barrier to corrosive elements of marine atmospheres. 1. 2. 3.
R.C. Tucker, “Thermal Spray Coatings”, ASM International, Vol. 5 (1996), p. 1446 G. Vourlias et al., High Temp. Mat. Proc., 9 (2005), p. 243. ASM handbook of metals, “Forms of Corrosion”,Vol.13 (1994), p. 97.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 705–706, DOI: 10.1007/978-3-540-85226-1_353, © Springer-Verlag Berlin Heidelberg 2008
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Figure 3. Cross Section SEM micrographs of NiCrBSi (a) and Al (b) coated steel after 1 month of corrosion in the salt spray chamber with inset chemical mapping images.
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Microscopical study of the influence of zinc addition on the structure of WO3 K. Nikolaidis, D. Chaliampalias, G. Vourlias, E. Pavlidou, G. Stergioudis Physics Department, Aristotle University of Thessaloniki, 54124 Thessaloniki, Greece [email protected] Keywords: zinc, Electron Microscopy, XRD, tungsten
Tungsten oxide is used in several electrical, optical and chemical applications. The addition of metal ions in its structure can improve further its advantageous properties. The WO3 structure contains lattice voids, in tunnel form, parallel to b-axis, where several metal ions can be inserted. This ion intake can be further motivated by temperature. By the continuous metal ion insertion, the structure of WO3 transforms gradually from monoclinic to tetragonical and final to cubic phase [1-3]. The aim of the present work is to examine the influence of zinc ion insertion, at several concentrations, into the WO3 structure. This is performed, initially, by mixing WO3 powder with zinc powder from 1 to 35%wt. Then the mixtures are enclosed in particular ampoules under 10-6 vacuum in order to facilitate the chemical reaction (xZn+WO3 ZnxWO3). The enclosed mixtures are then left for 2.5 days under 1066oC in a preheated furnace and after the necessary preparation, the specimens are examined with scanning electron microscopy (20kVolt JEOL 840A SEM equipped with an OXFORD ISIS 300 EDS analyzer) and XRD diffraction analysis. From SEM observations it is found that the crystal shape and morphology of ZnxWO3 powder changes with zinc concentration in the powder mixture. Particularly, at 1-5%wt. Zn, crystals form elevated size (10-30μm) grains (fig. 1a). From 6 to 15%wt. Zn, needle shaped crystal growth is favoured (fig. 1b, c). The length of the as formed needles varies from 10 to 200μm, especially near 15%wt. Zn. Finally at higher zinc concentrations, crystals form cubic shaped grains (fig. 1d) while several needle shaped grains are grown over them (fig. 2a, b). EDS analysis performed on mixtures containing 35%wt. zinc, revealed that cubic grains contain higher Zn amounts (fig.2c, area 1) while on needle grains (fig. 2c, area 2) higher W concentration is recorded (fig. 2d). Diffraction patterns from XRD analysis show that at low Zn concentrations (3%wt.) only the monoclinic phase of the system is formed. At 6%wt. Zn, it is the beginning of the system phase transformation as both monoclinic and tetragonical phases are recorded in the Zn0,06WO3 powder. At 15%wt. Zn, only the tetragonical phase is found while near 25%wt. Zn, some cubic phases are traced. Finally, at 35%wt. Zn, the cubic phase is predominant and only weak intensity peaks of the tetrgonical phase are recorded. 1. 2. 3.
A.R.Phani, M.Passacantando, L.Lozzi, S.Santucci, J. Mat. Sci., 35 (2000), p. 4879. Arne Magneli, Chemical Scripta, 26 1986, p. 535. J. Booth, T. Ekstrom, E. Iguchi and R. J. D. Tilley, J. Solid State Chemistry, 41 1982, p. 293.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 707–708, DOI: 10.1007/978-3-540-85226-1_354, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. SEM micrographs of ZnxWO3 powder with 3%wt. (a), 6%wt. (b), 15%wt. (c) and 25%wt.(d) zinc concentration.
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Microstructural Studies by Electron Microscopy Techniques of TiAlSiN Nanostructured Coatings V. Godinho1,2, T.C. Rojas1, M.C. Jimenez, M.P. Delplancke-Ogletree2, A. Fernández1 1. Instituto de Ciencia de Materiales de Sevilla CSIC-US Avenida Américo Vespucio 49, 41092 Sevilla, Spain 2. Université Libre de Bruxelles, Chemicals and Materials Department, Faculty of Applied Sciences, Avenue F.D. Roosevelt, 50 (CP165/163), 1050 Bruxelles, Belgium [email protected], [email protected] Keywords: coatings, TiAlSiN, nanocomposites, microsctructure
In the last decade the industry demands for hard coatings with enhanced thermodynamic stability has increased the interest in the study of superhard TiAlSiN nanocomposite coatings[1]. A full characterization of nanostructured TiAlSiN coatings at the microstructural and compositional level is needed to understand the good mechanical properties of these coatings[2,3]. In this work, the coatings have been deposited onto different substrates by reactive magnetron sputtering using two commercial TiAl (75/25 at. %) and Si (99.999 at. %) targets. The influence of experimental parameters such as power of the sputtering source (100, 250, 400 y 600 W) and appliance of a bias voltage of 25W during deposition has been correlated with the structure and composition of the coatings. The combination of different characterization techniques such as Scanning Electron Microscopy (SEM-FEG), Transmission Electron Microscopy (TEM), Selected Area electron diffraction (SAED) and Electron Energy Loss Spectroscopy (EELS) will be presented as a suitable methodology to characterize nanostructured coatings in general and the TiAlSiN coatings in particular. The SEM y TEM study show that the samples prepared with bias present very smooth surface and dense structures while the coatings prepared without bias show columnar structure resulting in a rough surface (Fig 1). This can be very important regarding the mechanical properties and final applications. An increase in the power supply increases the thickness of the film. TEM/SAED analysis revels that the samples prepared at lower power are amorphous. The increase of the power of the TiAl target increases the crystallinity of the samples. For the same power the non biased sample present crystal sizes bigger than the biased ones, (Fig. 2). The indexed diffraction pattern corresponds to the Ti3AlN cubic phase. The TiN cubic phase can also be present. EELS results also revealed the formation of TiAlN, and the presence of amorphous SiONx phases. Changes in N-edge shape can be observed when the power on the TiAl targed is increased. At higher TiAl target power the presence of TiN phase is also detected. 1.
S. Veprěk, S. Reiprich, Thins Solid Films 268 (1995)64-71
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 709–710, DOI: 10.1007/978-3-540-85226-1_355, © Springer-Verlag Berlin Heidelberg 2008
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S.Carvalho, L.Rebouta, A.Cavaleiro, L.A.Rocha, J.Gomes, E.Alves, Thin Solid Films 398399 (2001) 391-396 G.S.Kim, B.S.Kim, S.Y.Lee, J.H.Hahn, Thins Solid Films 506-507 (2006) 128-132. We acknowledge the financial support from the Spanish MEC (Proyect no. MAT2000401052 and MAT2007-66881-C02-01) and NOE EXCELL NMP3-CT-2005-515703
Figure 1. SEM micrographs in cross section for samples at 600W without (left) and 25W biased (middle). Cross section TEM detail of the sample at 600W with bias (right).
Figure 2. TEM micrographs and electron diffraction patterns for the samples at 400W. Indexed planes correspond to the Ti3AlN cubic phase. EELS N-K edge of the samples with 0 and 25W of sustrate bias
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Structural and interface studies of a nano-scale TiAlYN/CrN/alumina coating I.M. Ross1, W.M. Rainforth1, C. Strondl2, F. Papa2 and R. Tietema2 1. Department of Engineering Materials, University of Sheffield, Mappin Street, Sheffield S1 3JD, United Kingdom 2. Hauzer Techno Coating BV, Van Heemskerckweg 22, NL-5928 LL Venlo, Netherland [email protected] Keywords: TiAlYN/CrN, Alumina, Electron Energy Loss Spectroscopy
Protective nitride overlay coatings based on nano-scale multi-layer structures currently represent the state-of-the art in industrial coating technology [1]. In the present study, a TiAlYN/CrN multi-layer coating was deposited onto a γ-TiAl substrate using combined cathodic-arc/unbalanced magnetron sputtering. To further extend the oxidation and wear resistance of the coating a newly developed alumina top coat was subsequently deposited by pulsed DC sputtering. We have utilised conventional HREM imaging, electron diffraction and analytical STEM (EELS and EDS) to investigate the substrate/coating interface and structure of the alumina top coat. An overview of the coating architecture is provided in Fig. 1a. Detail of the interface between the nitride coating and the top coat is given in Fig. 1b. Selected area electron diffraction patterns revealed a diffraction pattern, indicative of a random nanocrystalline structure, which could be partially indexed to the cubic γ-Al2O3. Electron energy-loss studies of the near edge structure (ELNES) at the Al L2,3 and OK edge were investigated for the top coat (Figs. 2a and 2b respectively). The edge shape and peak positions reflect the local atomic environment and can act as a fingerprint for the local atomic coordination. By comparison with standard energy-loss spectra for both the Al L2,3 and O K a good agreement was achieved for the γ-Al2O3 phase, supporting the results of the electron diffraction studies [2]. Prior to deposition the substrate surface was pre-treated using a Cr ion etch. This has been shown to improve adhesion compared with conventional Ar+ plasma treatments. The interface between the γ-TiAl substrate and TiAlN base-layer, shown in Fig 3a, revealed evidence for localized epitaxy. Energy-loss spectra obtained across the interface, (Fig. 3b) show two important features: (i) Cr implantation at the substrate surface, consistent with the Cr ion etch and (ii) the presence of nitrogen to a depth of ~6nm within the surface region of the substrate. It is believed that the implantation process reduces lattice mismatch between the substrate and coating, reducing interface stresses hence resulting in improved adhesion. 1. 2.
C. Leyens, M. Peters, P.Eh. Hovsepian, D.B. Lewis, Q. Luo, W.D. Münz, Surf. Coat. Technol. 155 (2002), p. 103 I. Levin, A. Berner, C. Scheu, H. Muellejans, D.G. Brandon, Mikrochim. Acta. 15 (1998), p. 93
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3.
The Authors gratefully acknowledge the support of the European Commission (INNOVATIAL-Project NMP3-CT-2005-515844).
Figure 1. (a) Low magnification STEM bright-field (BF) image showing the substrate, base-layer, multi-layer and lower part of the alumina topcoat. (b) HREM image showing the interface between the multilayer coating and the alumina top coat.
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Figure 3. (a) HREM image of the interface between the γ-TiAl substrate and TiAlN base-layer and (b) corresponding EEL spectra across the interface: (i) substrate, (ii) interface and (iii) base-layer.
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An investigation of SiC-fiber coatings T. Toplišek, Z. Samardžija, G. Dražić, S. Kobe, S. Novak Department for Nanostructured Materials, Jožef Stefan Institute, Jamova cesta 39, 1000 Ljubljana, Slovenia [email protected] Keywords: SiC fibers, electron microscopy, AFM
Continuous SiC fibers are one of the high-strength fibers [1] that can be used for reinforcing a ceramic matrix composite. In our case this is in a SiC/SiC composite material. Because of its excellent mechanical properties, low induced radioactivity after neutron irradiation, high strength, its specific thermal conductivity and low porosity it can be used as a functional material in a future fusion reactor [2]. High-strength fibers are stiff and have a low density, and the elements are covalently bonded. The flexibility of the fibers is connected with the Young’s modulus and the diameter. Fibers with a smaller diameter have a higher flexibility and vice versa. That is why we can easily weave, bend and twist them into complicated forms. In this investigation we used amorphous SiC fibers (Nicalon) with a diameter of 15 μm and crystalline SiC fibers (Tyranno SA) with a diameter of 7 μm. It is known that the properties of the fiber/matrix interface play an important role in determining the mechanical and physical properties of ceramic matrix composites. To prevent a catastrophic failure of the composite material and to control the crack deflection at the interface [3] we protect the fibers with different thin layers, diamondlike carbon (DLC), CrC and WC, called the “interphases”, using magnetron sputtering. The fibers were investigated with scanning electron microscopy (SEM), transmission electron microscopy (TEM) and atomic force microscopy (AFM) before and after the coating. From Figure 1a we can see that the Nicalon SiC fibers have smooth surface and are without any visible defects. Transmission electron microscopy (TEM) analyses showed that the Nicalon SiC fibers consist of an amorphous matrix phase, which contains a small amount of nanocrystallites, 1–3 nm in size (Figure 1b). Their diffraction patterns can be indexed as cubic SiC. The average roughness (Ra) of the Nicalon fiber surface, determined with AFM, is 2.0 nm (Figure 1c). Figure 2a shows an SEM micrograph of the WC coating on the Nicalon SiC fiber. The thickness of the layer varies between a few nm and 0.5 μm, depending on the experimental conditions and the overlapping of the fibers during the deposition. The TEM analysis shows that in the case of the WC coating, a multilayer (sandwich structure) approach was used (Figure 2b). The first layer on the fiber surface is chromium and the second is tungsten. The role of these two intermediate layers is to ensure better cling and adhesion of the WC coating to the fiber surface. The third layer is tungsten carbide, with a thickness of 400–500 nm. The diffraction patterns show that the chromium layer is crystalline; on the other hand, the WC layer is amorphous. The chemical composition of the individual phases was determined with TEM using energy-
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dispersive x-ray spectroscopy (EDXS). The average roughness of the Nicalon fiber coated with WC is 7.3 nm (Figure 2c). 1. 2. 3. 4.
K. K. Chawla, Composite Materials: Science and Engineering (2nd ed.), Springer (1998). T. Taguchi, N. Igawa, R. Yamada and S. Jitsukawa, Journal of Physics and Chemistry of Solids 66 (2005), p. 576-580. S. Bertrand, C. Droillard, R. Pailler, X. Bourrat and R. Naslain, Journal of European Ceramic Society 20 (2000), p. 1-13. We kindly acknowledge Mrs. Medeja Gec for her help in TEM sample preparation and Dr. Peter Panjan from the Department of Thin films and Surfaces, Jožef Stefan Institute, Slovenia for his help in depositing the thin layers on the fiber surface. This work was performed under the contract P2-0084 and was financially supported by the Ministry of Higher Education, Science and Technology of the Republic of Slovenia and the European Commission within the Contract of Association Euratom FU06-CT-2004-00083.
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Figure 1. SEM (a), TEM (b) and AFM (c) micrographs of amorphous Nicalon SiC fiber. a
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Figure 2. SEM (a), TEM (b) and AFM (c) micrographs of Nicalon SiC fiber coated with WC.
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HRTEM-EELS study of atomic layer deposited thin rare earth oxide films for advanced microelectronic devices S. Schamm, P.E. Coulon and L. Calmels CEMES-CNRS, BP 94347, 31055 Toulouse cedex 04, France [email protected] Keywords: HRTEM, EELS, ELNES, Rare Earth Oxides, High Dielectric Constant, Gate Dielectrics, Atomic Layer Deposition
Replacing the SiO2 gate dielectric in advanced microelectronic devices to reduce leakage and improve gate capacitance is the subject of intensive research since nearly a decade. Several materials with a high dielectric constant (k) have been considered. But only few among them are eligible due to the very stringent requirements for the materials properties such as bandgap and band alignment to Si and chemical and thermal stability (up to 1000°C due to processing conditions). Particularly, because the required thickness of the dielectric films is today in the nanometer range, reliability becomes a major problem to solve due to instability issues. Thus different parameters of the multilayer dielectric stack such as (i) the roughness at the interfaces with the semiconductor substrate and the gate, (ii) the thickness of the unavoidable interfacial layer (IL) that forms between the Si substrate and the high-k film, (iii) the uniformity of chemical composition in the IL and within the high-k film must be controlled strictly [1]. Rare earth oxides (REO) are good candidate to consider for substituting SiO2 due to their high conduction band offsets and moderately high k values [2]. Among the REO series, La2O3 and Lu2O3, with empty and completely filled f shell respectively, are the most promising oxides. La2O3, which is stabilized in the hexagonal phase above 400 0C, features the highest k value (k=27). However, it is the most hygroscopic oxide of the REO series exhibiting the highest affinity for Si atoms. On the contrary, Lu2O3, despite a medium k value (k=12.5), is the least hygroscopic oxide among the above series. It is also less prone to Si atom diffusion than any other oxide of the same series. In this presentation, we propose to show how transmission electron microscopy performed in the high-resolution mode (HRTEM) "Figure 1" coupled with electron energy loss spectroscopy (EELS) can afford valuable information in critical parameters for advanced microelectronic devices [3] particularly concerning the structure and chemical composition of the IL and of the film atop. This will be illustrated with the case of atomic layer deposited (ALD) thin films based on La or Lu oxides prepared on a Si substrate. Modelisation of O-K fine structures in compounds of the La-O-H-Si "Figure 2" and Lu-O-H-Si systems will be proposed for phase identification. The (FP6-2004-IST-NMP-3) Project REALISE (2006-9) is acknowledged for funding. The authors thank G. Scarel (MDM-Agrate, Milan) for fruitful discussions and Lu-based films growth and H.L. Lu (MDM-Agrate, Milan) for La-based films growth.
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1. 2. 3.
H. Wong, H. Iwai, Microelectronic Engineering 83 (2006), p.1867. G. Scarel, A. Svane, and M. Fanciulli, in "Rare earth oxide thin films: growth, characterization, and applications", eds. M. Fanciulli and G. Scarel, (Springer-Verlag), Topics in Applied Physics 106 (2006) p1. S. Schamm, G. Scarel and M. Fanciulli, in "Rare earth oxide thin films: growth, characterization, and applications", eds. M. Fanciulli and G. Scarel, (Springer-Verlag), Topics in Applied Physics 106 (2006) p153.
Figure 1. High-resolution image of a ALD grown La-based film showing the stack: Si substrate, amorphous interfacial layer and nanocrystalline high-k layer.
Figure 2. Oxygen K-edge fine structure for cubic and hexagonal La2O3 calculated with the multiple scattering DFT-based code FEFF8.
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New Fullerene like materials for tribological applications: TEM and EELS study Virginie Serin1, Nathalie Brun2, Ch. Colliex2 1. CEMES-CNRS, 29 rue Jeanne Marvig, 31055 TOULOUSE, France 2. LPS-CNRS, Université Paris Sud, 91405 ORSAY cedex France [email protected] Keywords: EELS, Materials Science
Fullerene-like materials can be defined as materials made up of two-dimensional lamellar basic units of nanoscale dimensions with closed and curved shapes. They can grow in the solid phase into three-dimensional networks or architectures displaying a wide range of morphologies. The basic origin of these new phases lies in the role of dangling bonds at the periphery of sheets of such small sizes, which destabilise planar structures and induce closing. The atomic-scale structure of these materials involving strong covalent inter-atomic bonding and non-compact space filling, make them strong candidates for breakthrough developments, of potentially high interest in many industrial fields. R. Tenne et al first discovered this new class of three-dimensional inorganic nanostructures [1]. They found that certain inorganic compounds that normally occur as large flat platelets can be synthesized into much smaller nano-spheres. Due to their size, shape, chemistry, and structure, materials based on these nanoparticles have special properties compared to materials based on conventionally sized constituents of the same composition. This makes them attractive for many commercial applications. Up until the discovery by the Weizmann group it was thought that fullerenes could only be made with carbon atoms. They were the first to discover that certain inorganic (i.e., noncarbon) materials could also be formed into fullerene-like structures, hence the name inorganic fullerene-like materials, or IFLM nanoparticles. In the present paper we report the study of samples developed with the overall objective to provide industry with radically new composite coating systems. The application of the composite coatings will be for surfaces and lubricants, in order to significantly reduce and control friction and wear in tribological contacts. The ultimate aim is to reduce friction as well as to extend operational life, reduce maintenance requirements and reduce the environmental impact of a wide range of mechanical elements for the aerospace, automotive, power generation (energy) and manufacturing industrial sectors [1, 2,3]. The studied samples have been prepared by different methods. There are WS2 nanoparticles incorporated in different types of coatings: Cr2O3, siloxane, TiN, Ni:P or DLC (see figure 1), or (ii) MoS2 nanoparticles in situ formed within a DLC or TiN matrix (see figure 2). In this presentation we will try to connect structural, chemical and bonding information to the tribological properties of the samples.
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The authors are pleased to acknowledge the financial support of the EC (FOREMOST, Contract no.: FP6-515840) 1. 2. 3.
L. Rapaport, Y. Feldman, M. Homyonfer, H. Cohen, J. Sloan, J.L. Huchison, R. Tenne, Inorganic-like fullerene-like material as additives to lubricants; structure-function relationship, Wear, 225-229 (1999) 975-982 M. Choowalla and G.A.J Amaratunga, Thin films of fullerene-like MoS2 nanoparticles with ultra-low friction and wear, Nature, 407 (2000) 164 – 167. L. Rapaport, V Leshinsky, M. Lovovsky, I. Lapsker, Y. Volovik, Y. Feldman, R. PopovitzBiro, R. Tenne, Superior tribological properties of powder materials with solid lubricant nanoparticles, Wear 255, (2003) 794-800.
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Figure 1: Left: HRTEM of a WS2 single nanoparticle, Middle: WS2 nanoparticle in Cr2O3, Right: mapping of an interface WS2@siloxane: W (red), O (green), Si (blue).
10 nm Figure 2: MoSX fringes in situ formed in two different coatings.
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Phase determination of nanocrystalline Al-Cr-O coatings by analytical TEM J. Thomas1, J. Ramm2, B. Arnold1, B. Widrig2, T. Gemming1 1. IFW Dresden, P.O. Box 27 01 16, 01171 Dresden, Germany 2. OC Oerlikon Balzers AG, Iramali 18, LI-9496 Balzers, Liechtenstein [email protected] Keywords: coatings, nanocrystallinity, analytical TEM
Coatings are an important possibility to improve the performance of tools. Hardness, brittleness, layer stress, coefficient of friction, and heat-resistance are essential properties of such coatings. However, some properties like hardness and brittleness are in opposition to each other. One of the possibilities to overcome such oppositions is the deposition of thin nanocrystalline coating layers in which the higher contribution of surface energy to the thermodynamical potential can lead to completely new compositions and phases. The characterisation of these coatings needs a method sensitive for thin layers with high spatial resolution for imaging, phase determination and chemical analysis. These challenges are met by the analytical transmission electron microscopy (TEM). The investigated Al-Cr-O layers were deposited by pulse enhanced electron emission (P3e™), a PVD technology for the synthesis of metal oxides utilizing cathodic arc evaporation in which the arc current can be pulsed. For the TEM characterisation on cross sections, electron-transparent lamellas were cut by focused ion beam (FIB) lift-off technique. For the characterisation a TEM/STEM Tecnai F30 superTWIN equipped with EDXS detector and imaging filter GIF200 has been used. Figure 1 shows a TEM lamella before the lift-off step. The described layer system is typical for the investigated samples. The morphology, the phase formation and the locally resolved chemical composition of the coating layer are in the focus of interest. Especially the determination of the phases plays an important role. Because of the nanosized grains, an overlay of different diffraction patterns has to be expected. Beside the determination of the existing phases, their quantities are also of interest. To solve this problem, a combination of electron diffraction and EDXS measurements has been used. For given phases, the EDXS results issue a system of equations for the phase mixture. The solution has to be confirmed by the quantitative calculation of the diffraction intensities of an overlay of the phases. An example for the results is given in Figure 2. The cross section image 2a) with the inserted diffraction patterns show the nanocrystalline morphology of the coating layer. This is confirmed by the HRTEM image 2b) that demonstrates grain sizes in the order of 5nm. The radial density distributions (Figure 2c) extracted from the patterns A and B indicate that there is no gradient of phase formation within the coating. The diffraction intensity curve can be explained quantitatively by assuming a mixture of 77 at-% Al2O3, 14 at-% Cr2O3 and 9 at-% Al which is in sufficient agreement with EDXS results.
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Figure 1. Scanning electron microscopic image of a TEM lamella before lift-off with description of the typical layer system.
Figure 2. Morphology and phase formation within the coating layer. a) TEM cross section micrograph with diffraction patterns of positions A and B. b) HRTEM image with fourier-filtered insets. c) radial density distributions of patterns A and B. d) best fit of the calculated overlay of phases (see text).
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Bioinspired synthesis of nanostructures based on S-layer lattices D. Pum, N. Ilk and U.B. Sleytr Center for Nanobiotechnology, University of Natural Resources and Applied Life Sciences, Gregor Mendelstrasse 33, A-1180 Vienna, Austria [email protected] Keywords: protein assembly, nanoparticle arrays, biomineralization
One of the most challenging research areas is currently found at the interface between life and non-life sciences. Self-assembly of molecular building blocks and bottom-up fabrication of supramolecular architectures has become an important scientific and engineering discipline crossing the boundaries of several established fields. In particular, the bioinspired synthesis of nanostructured materials with novel structural, mechanical, electrical, optical or catalytic properties has only become possible due to the convergence of physics, chemistry, and biology. Monomolecular bacterial surface layer proteins (S-layer proteins) are versatile assembly systems providing a structural basis for a complete supramolecular construction kit involving all major species of biological molecules (proteins, lipids, glycans, and nucleic acids) [1]. S-layers are the most commonly observed cell surface structures in prokaryotic organisms (bacteria and archaea) [1]. They are composed of a single protein or glycoprotein species (Mw = 40 to 200 kDa) and exhibit either oblique, square or hexagonal lattice symmetry with unit cell dimensions in the range of 3 to 30 nm. Slayers are generally 5 to 10 nm thick. They represent highly porous protein meshworks with pores of uniform size and morphology in the 2 to 8 nm range. One of the key features of isolated S-layer proteins is their intrinsic tendency to self-assemble into monomolecular arrays in suspension, at solid supports (e.g. silicon wafers), at the airwater interface, at lipid films, liposomes, and nanocapsules. This presentation reviews the reassembly of native and genetically functionalized Slayer proteins and, in particular, their use as matrices for the templated assembly of molecules and nanoparticles into ordered arrays. In a wet chemical approach which was derived from fine grain mineralization in bacteria [2] self-assembled S-layer structures were exposed to metal-salt solutions followed by slow reaction with a reducing agent or by electron irradiation in an electron microscope [3-6]. The latter approach is technologically important since it allows the definition of areas where nanoparticles are formed [4,6]. Based on the work on binding biomolecules, such as enzymes or antibodies, it has also been demonstrated that metallic and semiconducting nanoparticles can be bound in regular arrangements (Figure 1) [7-9]. In both approaches nanoparticle superlattices were formed according to the lattice spacing and symmetry of the underlying S-layer. Although native S-layers have clearly demonstrated the presence and availability of functional sites for the precipitation of metal ions or binding of nanoparticles, a much more controlled and specific way of making highly ordered
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nanoparticle arrays uses genetic approaches for the construction of chimeric S-layer fusion proteins [10] incorporating unique polypeptides which have demonstrated to be responsible for biomineralization processes. Such chimeric S-layer protein lattices can be used as self-assembled nanopatterned affinity matrices capable to arrange molecules and nanoparticles into ordered arrays on surfaces [11]. 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11.
U.B. Sleytr, E.M. Egelseer, N. Ilk, D. Pum and B. Schuster. FEBS J. 274 (2007), p.323. S. Douglas and T.J. Beveridge. FEMS Microbiol. Ecol. 26 (1998), p.79. W. Shenton, D. Pum, U.B. Sleytr and S. Mann. Nature 389 (1997), p.585. S. Dieluweit, D. Pum and U. B. Sleytr. Supramol. Sci. 5, (1998), p.15. M. Mertig, R. Kirsch, W. Pompe and H. Engelhardt. Eur. Phys. J. D 9 (1999), p.45. R. Wahl, M. Mertig, J. Raff, S. Selenska-Pobell and W. Pompe. Adv. Mat. Sci. Technol. 13 (2001), p.736. E. Györvary, A. Schroedter, D.V. Talapin, H. Weller, D. Pum and U.B. Sleytr. J. Nanosci. Nanotech. 4 (2004), p.115. S. R. Hall, W. Shenton, H. Engelhardt and S. Mann. Chem. Phys. Phys. Chem. 3 (2001), p.184. M. Bergkvist, S.S. Mark, X. Yang, E.R. Angert and C.A. Batt. J. Phys. Chem. B 108 (2004) p.8241. D. Moll, C. Huber, B. Schlegel, D. Pum, U. B. Sleytr and M. Sára. Proc. Natl. Acad. Sci. USA 99 (2002), p.14646. We kindly acknowledge the partial support of AFOSR (Project FA9550-06-1-0208).
Figure 1. Schematic drawing of the different possibilities of binding nanoparticles on Slayers: (a) electrostatic binding, (b) covalent binding, and (c) using highly specific functional domains in S-layer fusion proteins.
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Direct Imaging of Carbon Nanoparticles inside Human Cells Alexandra E. Porter1, Crystal Cheng2, Mhairi Gass3, Karin Muller4, Jeremy Skepper4, Paul Midgley5, Mark Welland2 1. Department of Materials Science, Imperial College London, U.K., 2. The Nanoscience Centre, University of Cambridge, 11 J. J. Thompson Avenue, Cambridge CB3 OFF, UK, 3. UK SuperSTEM, Daresbury Laboratory, Daresbury, Cheshire WA4 4AD, UK, 4. Multiimaging Centre, Dept. of Anatomy, University of Cambridge, Downing Street, Cambridge CB2 3DY, UK, 5. Dept. of Materials Science and Metallurgy, University of Cambridge, Pembroke Street, Cambridge CB2 3QZ, UK [email protected] Keywords: carbon nanoparticles, toxicity, energy-filtered TEM
The application of nanotechnology in disciplines as varied as medicine and electronics is advancing rapidly with carbon nanoparticles (CNPs) such as fullerenes (C60) and nanotubes at the forefront. However, a lack in understanding of the interaction of such small structures with cellular material has resulted in concerns over their impact on human health [1-3] and since the individual structures have a diameter of ~1 nm they are potentially small enough to penetrate through ion channels or diffuse through pores in the nuclear membrane. Assessing their toxicity is imperative. In response to these concerns there has been an increase in the number of papers addressing the toxicity of carbon nanoparticles over the last few years but much of this data appears contradictory [1-6]. It is therefore essential to understand how the human body interacts with CNPs and more specifically to elucidate pathways by which CNPs enter the cell and their distribution within. SWNTs have been shown to be acutely toxic [1-3] in a variety of cells but the direct observation of cellular uptake of SWNTs has not been demonstrated previously due to difficulties in discriminating carbon-based nanotubes from carbon-rich cell structures. We exposed human monocyte derived macrophage cells to SWNTs for 2 and 4 days and showed that it is possible to map the location of intracellular SWNTs using energy filtered TEM (EFTEM) and confocal microscopy (Figure 1). We successfully imaged individual SWNTs within lysosomes and also crossing cell membranes (Figure 1, 2). We demonstrate two possible pathways of entry of SWNT into cells: energy-dependent phagocytosis or endocytosis and passive diffusion across lipid bilayers. SWNTs were found primarily within lysosomes, but more significantly within the cytoplasm and the nucleus (Figure 2). Uptake to these sites implies they may interact with intracellular proteins, organelles and DNA which would greatly enhance their toxic potential. SWNTs also fused with the plasma membrane where they have been shown to cause cell damage via lipid peroxidation and oxidative stress [2,3]. Localization of SWNTs at these sites was correlated to an increase in cell death in both cell viability assays and TEM analysis. To conclude we show that direct imaging of SWNTs within cells is
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achievable and is essential to complement cytotoxicity assays to understand localized effects of SWNTs and establishment of their potential toxicity. 1
Shvedova AA, Castranona V, Kisin ER, Schwegler-Berry D, Murray AR, Gandelsman VZ and Baron P, J. Toxicol. Environ. Health A, 66 2003 p. 1909. Manna SK, Sarkar S, Bar J, Wise K, Barrrera EV, Jejelowo O, Rice-Ficht A and Ramesh G, Nano Lett. 5(9) (2005) p. 1676. Chui D, Tian F, Ozkan CS, Want M, Gao H, Toxicol. Lett. 155 (2005) 155 p.73 Worle-Knirsch JM, Pulskamp K and Krug HF, Nano Lett. 6(6), 2006 p.1261 Sayes CM, Liang F, Hudson JL, Mendez J, Guo W, Beach JM, Moore VC, Doyle CD, West JL,Billups WE, Ausman KD and Colvin VL, Toxicol. Lett. 161(2) (2006) p135 We kindly acknowledge the EPSRC and The Oppenheimer Research Fellowship for funding.
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Figure 1. Cells exposed to SWNTs for 2 days. a) HAADF-STEM image of SWNTs within a lysosome; b) the low-loss EELS spectra from the cell alone and SWNTs in the cell show a small shift in plasmon energy; c) the iron edge extracted from the low loss spectrum which is integrated, the resulting iron map is shown in d) separating out the contribution from SWNTs and iron. Mapping the shift in plasmon energy results in e) where SWNTs can be identified for high resolution BF imaging f); g) another example of SWNTs end on in the cytoplasm, ends highlighted in yellow. a b
Figure 2. a) BF-STEM image of SWNTs translocating across the lipid bilayer into the neighboring cytoplasm. b) Confocal microscope image of HMM exposed to bundles of AgI@SWNT at 3 days confirming inclusion of SWNT bundles inside the nucleus (blue).
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Micro- and Nano-Textured Surfaces on Ti-Implants Made by Various Methods U. Beck1, R. Lange1 and H.-G. Neumann2 1. DOT GmbH, D-18059 Rostock, Germany 2. University of Rostock, IEF, Inst. GS, D-18119 Warnemünde, Germany [email protected] Keywords: Titanium, surface, nano-structure, electrolytic plasma, etching
Titanium and Titanium alloys have been used widely and successfully for various types of load bearing bone-anchored implants. It has revealed that the surface topography has a great importance thereby. Also the properties of oxide films or calcium phosphate (CaP) coating covering these implant surfaces are important for the osseointegration [1-5]. One interesting way to produce such oxide films are micro-arc or micro-plasma processes [6]. CaP is well-known as a substance which is highly similar to body-own material. The calcium phosphates can provided as brushite and hydroxyapatite, the minerals that take part in the formation of bone tissue. CaP coatings can be fully resorbed by the human body. The aim of our investigations here was to texture the implant surface at the microand nano-scale and to test various methods and processes with respect to their desired surface effects. In our project we used chemical etching to remove corundum particles remaining on the surface after blasting process - because of their possibly negative influence on osteointegration [7] - and various electric processes like electrolytic plasma and also cathodic processes to create new surface structures. As sample-base we used only cp Titanium. The CaP coating was precipitated electrochemically corresponding to the method described elsewhere [4, 8]. For etching and micro-plasma treatments all samples firstly were blasted with corundum (Al2O3) particles (3.5 bar) for producing of a coarse surface structure. Then all samples were chemically etched to remove the remaining corundum particles from blasting process and to cause a basic fine structure. Then the samples were treated anodic by a so-called electrolytic micro-plasma process which takes place at the interface between a solid (metals like Al or Ti) and a liquid electrode. The results of these treatments are shown in the following figures. The textured surfaces and coatings obtained were examined by SEM (FESEM SUPRA 25, Zeiss, Germany and S360, LEO, Germany) and EDX (Quantax, Bruker, USA).
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Figure 1., 2. Ti surfaces - chemically etched after corundum blasting (left side) and after micro-plasma process (right side); x 5000
Figure 3., 4. Various CaP structures on cp Titanium (left side x 5000, right side x 10000) 1. 2. 3. 4. 5. 6. 7. 8.
Y.T. Sul et al.: Int. J. Oral. Maxillofac. Implants, May-Jun. 2005, 20(3) p. 349-59 Y.T. Sul et al.: Biomaterials, Nov. 2005, 26(33) p. 6720-30 Y.T. Sul et al.: Biomaterials, Jan. 2002, 23(2) p. 491-501 P. Zeggel: Intern. Mag. Oral Implant., 1(2000) p. 52-57 P. Becker et al.: Key Engg. Mat., 218-220(2002) p. 63-656 P. Huang et al.: J. Biomed. Mater. Res., Aug 15 2004, 70B(2) p.187-90 J. Teller et al.: Patent DE 4431862 (1994). H.-G. Neumann, P. Zeggel and P. Becker: Patent WO 02/05862
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Determination of the biocompatibility of biomaterials by scanning electron microscopy (SEM) M. Bovi1,2, N. Gassler1,2 and B. Hermanns-Sachweh1,2 1. Electron microscopic facility of the medical faculty of the RWTH Aachen, Germany 2. Institute of Pathology, RWTH Aachen, Germany [email protected] Keywords: sem, eds, low vacuum sem
Scanning electron microscopy is an excellent method to determine the biocompatibility of biomaterials. With the possibility of showing the surface of biomaterials and measuring the surface properties of the material of interest like surface roughness, fins on bone implants or homogeneity of surface layers this method has also the potential to show the interface between biomaterial and biological layers. This article gives an overview of using the sem method for the examination of biomaterials in contact with cells and tissues. Hemocompatibility of biomaterials is tested by scanning electron microscopy in examining the reaction of thrombocytes on different surfaces. To show the interface between biomaterial and tissue a bone implant with the surrounding bone tissue is embedded in epoxy resin and polished until the cross section between implant and tissue is visible. Then with backscattered electron detector in low vacuum mode (no conductive layer necessary) the interface between implant and bone tissue is clearly seen “Figure 1”. The energy dispersive analysis of x-rays provides additional information of the elemental composition of the surface including the examination of extrinsic particles or demonstration of calcification on heart valves “Figure 2”. An elemental mapping shows where the different elements are located “Figure 3”. With the possibility of performing variable pressure scanning electron microscopy (VPSEM) there is no need of applying a conductive layer. But beside the backscatter electron image a real secondary electron image is still available [1]. These are some examples why scanning electron microscopy is a helpful tool for the examination of biomaterials in contact with cells and tissues and a testing device for biocompatibility. 1.
BJ. Griffin, Methods Mol Biol 369 (2007), p. 467-95
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Figure 1. High-resolution micrograph of the interface between implant and bone (white: implant, grey: bone tissue, black: resin)
Figure 2. High-resolution micrograph and EDX-spectrum of a heart valve with calcification
Figure 3. Crossection of aortic vessel with elemental mapping ( blue: C, yellow: Ca)
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Quantitative evaluation of the long-term marginal behaviour of filling restorations of human teeth using three-dimensional scanning electron microscopy W. Dietz1, S. Nietzsche1, R. Montag2, P. Gaengler3 and I. Hoyer2 1. Centre of Electron Microscopy, University of Jena, 07743 Jena, Germany 2. Dept. Conservative Dentistry, University of Jena, 07743 Jena, Germany 3. Dept.Conservative Dentistry, University of Witten/Herdecke, 58448 Witten,Germany [email protected] Keywords: quantitative electron microscopy, biomaterials, dental fillings
Clinical and micromorphological evaluations of the marginal behaviour of dental fillings are most important in restorative dentistry. Their scanning electron microscopical (SEM) investigation is well established [1,2]. However, the accurate quantification of margin changes (negative ledge, filling material loss) is not practicable by using common secondary-electron imaging. The results are affected by the subjective assessment of the examiner [3]. The aim of this study was to enhance the objectivity of the evaluation of dental filling materials by means of quantitative three-dimensional scanning electron microscopy (3-D SEM) [4]. Of 194 originally placed fillings (Visio Molar radiopaque, ESPE, Germany; occlusial/ approximal), 32 restorations were SEM investigated and 10 of these by 3-D SEM. Immediately after application (baseline), after 1, 5, 10 and 15 years, two-stage-replicas were taken from the patients´ teeth, gold sputtered and examined in the SEM Philips 515 using a four-quadrant backscattered electron (4-Q BE) imaging system (point electronic, Germany). 3-D reconstructions and quantifications (9 relief profiles of the very same places of each specimen) were computed using the topographical software MeX (Alicona, Austria) (Figure 1). The marginal ledges (enamel>filling) of the baseline varied from +7.7µm (marginal excess) to -20.3µm (negative ledge), average value -2.0µm (Figure 2). After 5 or 10 years at the latest, the maximum loss of the filling material above the interface was estimated between -7.2 and -48.4µm, depending on the baseline situation, average value: -27.0µm after 10 years. A fifteen-year use of most of the restorations showed reduced negative ledges of between -9.0 and -46.8, average value: –22.0µm. These investigations proved an increasing marginal filling material loss during the first five years (Figure 1a, b) causing negative enamel-filling ledges. Between 5 and 10 years, the margin near abrasion of enamel and filling was nearly the same (Figure 2). Between 10 and 15 years, the lowering of the negative ledges (Fig 1c, 2) was a result of the preponderance of the enamel abrasion. The quantitative long-term 3-D SEM evaluation of Visio Molar filling restorations of human teeth proved that the margin near loss of the filling material is limited and does not importantly affect the longevity of these restorations nor the secondary caries at the hard tissue-filling interface. This method is a valuable addition to the clinical and usual micromorphological evaluations. S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 729–730, DOI: 10.1007/978-3-540-85226-1_365, © Springer-Verlag Berlin Heidelberg 2008
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J.F. Roulet, B. Salchow and M. Wald, Dent. Mater. 7 (1991), 44. P. Gaengler, I. Hoyer, R. Montag and P. Gaebler, J. Oral Rehabil. 231 (2004), 991 R. Stoll, M. Gente, M. Palichleb and V. Stachniss, Dental Materials 23 (2007), 145 W. Dietz, S .Meineber, U. Kraft, I. Hoyer and E. Glockmann in Proceedings of the Microscopy Conference 2005, Davos Switzerland 28.8.-2.9.2005, ISSN1019-6447
100µm Pl
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E F
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1c Figure 1. Profile reconstruction of an enamel-filling (Visio Molar) margin. 1 (1a), 5 (1b) and 15 (1c) years after application. Pl: profile-line, E: enamel, F: filling.
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Figure 2. Long-term marginal filling material loss of dental fillings (Visio Molar): development of negative enamel-filling ledges (µm) over 15 years.
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The analysis of Si doped hydroxyapatite coatings using FIBSEM, TEM and RHEED H.K. Edwards1, S. Coe1, T. Tao1, M.W. Fay2, C.A. Scotchford1, D.M. Grant1 and P.D. Brown1 1. School of Mechanical, Materials and Manufacturing Engineering, 2. University of Nottingham Nanotechnology and Nanoscience Centre, both at the University of Nottingham, University Park, Nottingham, NG7 2RD, UK. [email protected] Keywords: FIBSEM, RHEED, EELS, EFTEM, hydroxyapatite, biomaterial
Silicon doping has been found to enhance the in vitro and in vivo bioactivity of hydroxyapatite (HA) [1,2], a bioceramic widely used in orthopaedic implants. The atomic arrangement of the Si within the HA structure and the exact role of the dopant on bioactivity are not completely understood [3]. Therefore, the effect of Si dopant levels on the HA thin film structure and chemistry has been examined using the combined techniques of transmission electron microscopy (TEM), selected area electron diffraction (SAED), scanning TEM (STEM), electron energy loss spectroscopy (EELS), energy filtered TEM (EFTEM) and energy dispersive X-ray (EDX) mapping, along with complementary scanning electron microscopy (SEM), X-ray diffraction (XRD) and reflection high energy electron diffraction (RHEED). Si-HA/Ti samples containing 0, 1.8, 4.2 and 13.4 wt.% Si were prepared by plasma assisted RF sputtering of HA (~200 nm thickness) using a multiple target unbalanaced magnetron configuration onto 10 mm by 1 mm Ti discs, followed by heat treatment at 600 ºC under flowing Ar. Due to the composite nature of the Si-HA/Ti samples, focused ion beam scanning electron microscopy (FIBSEM) has been employed for the preparation of cross sectional TEM samples. Specimens were prepared using an FEI Quanta 200 3D FIBSEM fitted with a Quorum cryo-transfer unit, an Omniprobe micromanipulator and an INCA Oxford Instruments EDX analysis system. Following deposition of a protective W coating, FIBSEM milling for lift-out was carried out using sequential ion beam currents of 20 nA down to 30 pA. Electron transparent specimens were inspected using a JEOL 2100f TEM and a liquid nitrogen cooled sample holder. Reference TEM and SAED results found that the pure HA thin film (0 wt % Si) had a polycrystalline structure (Figure 1). RHEED and XRD results indicated that a high concentration of Si destabilises the HA structure, causing the coatings to exhibit smaller HA crystallites and become increasingly amorphous. Preliminary STEM-EELS investigations showed a homogeneous Si, Ca and P dispersion throughout a 13.4 wt % Si - HA thin film on Ti, indicating that Si does not preferentially segregate to HA grain boundaries (Figure 2).
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 731–732, DOI: 10.1007/978-3-540-85226-1_366, © Springer-Verlag Berlin Heidelberg 2008
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I. Gibson, S. Best and W. Bonfield, J. Biomedical Materials Research, 44 (1999) p. 422. I. Gibson et al., Key Engineering Materials, 218 (2002) p. 203. D. Arcos et al., J. Biomedical Materials Research, 78A (2006) p. 762. This work was supported by the EPSRC under grant EP/E015379/1. The authors kindly acknowledge Joanne Hampshire at Teer Coatings Ltd. for the deposition of the HA thin films.
Figure 1. Bright field TEM image and SAED patterns of a pure (0 wt % Si) HA thin film on Ti and the protective W coating laid down during FIB preparation. The Ti and W coating were found to be crystalline while the HA thin film was polycrystalline.
Figure 2. Elementally sensitive maps derived from STEM-EELS analysis of a 13.4 wt % Si - HA thin film on Ti, showing a homogeneous distribution of Si, Ca and P within this coating on the 10 nm scale.
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Electron microscopic investigations of the polymer/mineral composite material nacre K. Gries1,2, R. Kröger3, C. Kübel4, M. Fritz1 and A. Rosenauer2 1. University of Bremen, Institute of Biophysics, Otto-Hahn-Allee 1, 28359 Bremen, Germany 2. University of Bremen, Institute of Solid State Physics, Otto-Hahn-Allee 1, 28359 Bremen, Germany 3. University of York, Department of Physics, Heslington, York YO10 5DD, United Kingdom 4. Fraunhofer Institute for Manufacturing Technology and Applied Materials Research (IFAM), Wiener Straße 12, 28359 Bremen, Germany [email protected] Keywords: nacre, TEM, mineral bridge
Nacre of the mollusc Haliotis laevigata was investigated using transmission electron microscopy (TEM). Nacre, the inner iridescent layer of mollusc shells, is a typical example for a material which is formed by biomineralization processes. It is composed of the CaCO3-polymorph aragonite and a small amount of about 5wt% organic matter [1]. The polygonal shaped aragonite platelets show a width which ranges from 5μm to 10μm and a thickness of about 500nm [2]. They are laterally arranged in layers and vertically in stacks. The aragonite crystal structure can be described by the space group Pmcn 62 with the c-axis perpendicular to the face of the platelets and thus to the surface of the shell [3]. The platelets are separated by layers of organic material, the organic matrix. This arrangement resembles a brick and mortar like structure. The combination of stiff crystalline material and soft organic material, as well as the layered structure increase the fracture toughness in comparison with pure (geological) aragonite. In TEM investigations structures which are located within the organic layer between stacked aragonite platelets were observable (marked in Figure 1 by a white arrow). Electron tomography investigations showed that these structures connect the stacked platelets. High resolution TEM allowed a detailed analysis of these structures, the socalled mineral bridges, and revealed that they consist of aragonite which exhibits a constant crystallographic orientation. During the process of growing the orientation of the aragonite might be transferred through the organic matrix via the mineral bridges. To check if stacked platelets exhibit a similar orientation, the tilt of these aragonite platelets against each other was determined from selected area diffraction patterns. It could be shown that the platelets show small relative tilt angles of <4° and that they exhibit a similar crystallographic orientation over a distance of about 15 platelets. Within the platelets facetted nanopores with diameters between typically 5nm and 15nm are observable (marked in Figure 1 by a black arrow). Only few nanopores have a larger diameter of about 30nm. Using electron tomography the position and the form of these pores could be determined. Electron energy loss spectroscopy (EELS) and energy
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dispersive X-ray (EDX) spectroscopy show that an increased amount of carbon and therefore possibly organic material is located within the nanopores. 1. 2. 3.
A. P. Jackson, J. F. V. Vincent and R. M. Turner, Proc. R. Soc. Lond. B 234 (1988), p. 415. M. Fritz, A. M. Belcher, M. Radmacher, D. A. Walters, P. K. Hansma, G. D. Stucky, D. E. Morse and S. Mann, Nature 371 (1994), p. 49. S. Mann in “Biomineralization: Principles and Concepts in Bioinorganic Materials Chemistry”, (Oxford University Press, New York) (2001), p. 118.
Figure 1. TEM image of a cross-sectional nacre specimen. The white arrow marks a mineral bridge that connects two stacked aragonite platelets. The black arrow marks a nanopore.
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Studies on the microstructure of fresh-cut melon I. Hernando, L. Alandes, A. Quiles, I. Pérez-Munuera Departamento de Tecnología de Alimentos, Universidad Politécnica de Valencia, Camino de Vera s/n 46022 Valencia, Spain. [email protected] Keywords: microstructure, melon, fresh-cut
Fresh-cut fruit is a segment of food industry that is developing fast due to the fact that it is a convenient food and has a fresh-like quality. However, fresh-cut fruits are still under research because of the problems involved in preserving their quality during prolonged periods. In order to improve their texture, postharvest calcium solution application has been used [1]. Calcium maintains the cell wall structure in fruits by interacting with pectic acid in the cell wall to form calcium pectate which firms molecular bonding between constituents of cell wall [2]. The aim of this work is to study the effect of calcium lactate on the microstructure of fresh-cut Piel de Sapo melon during three weeks of storage at 4ºC. A batch of melon cubes (15 mm side) was submerged in a 8 g·L-1 calcium lactate solution used as firming agent. Malic acid (25 g·L-1) was used as antimicrobial in order to prevent microbial growth during the study. The treatment was carried out at 10ºC for 1 min and the samples stored in polyethylene bags at 4ºC during 3 weeks. Another batch of melon cubes was only treated with antimicrobial, stored in the conditions described previously and considered as a control. The microstructure was studied by Low Temperature Scanning Electron Microscopy (Cryo-SEM) and Light Microscopy (LM). In the Cryo-SEM technique a JSM-5410 SEM microscope (Jeol, Tokyo, Japan) was used with a Cryo CT-1500C unit (Oxford Instruments, Witney, UK). The sample was frozen in slush nitrogen (T ≤ -210ºC), fractured, etched (-90ºC) and gold-coated (2mbar, 2 mA) for imaging, working at 15 kV and -130ºC. In the LM technique the samples were fixed with a 25 g·L-1 glutaraldehyde solution, postfixed with a 20 g·L-1 OsO4 solution, dehydrated in ethanol series, contrasted in 40 g·L-1 uranyl acetate and embedded in epoxy resin. The samples were cut using a Reichter Jung ultramicrotome (Leila, Barcelona, Spain), stained with toluidine blue and examined in a Nikkon Eclipse E800 light microscope (Izasa, Barcelona, Spain). After 1 week of storage at 4ºC, the control melon (Figure 1A) showed the intercellular spaces filled with solutes, while the calcium treated melon (Figure 1B) showed empty intercellular spaces. Moreover, closely bonded cells and intact cellular walls can be observed in the calcium treated melon. The cell walls of the control melon are thin, even some of them are broken and the plasmalemma starts retracting to the center of the cell (Figure 1C); in the calcium treated melon (Figure 1D) the plasmalemma remains close to the cell wall. After 3 weeks of storage at 4ºC, the cell degradation in the control melon parenchyma is higher (Figure 2A) than that of calcium treated melon (Figure 2B). The cell walls of the control melon are slightly stained due to S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 735–736, DOI: 10.1007/978-3-540-85226-1_368, © Springer-Verlag Berlin Heidelberg 2008
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the loss of fibrilar packing and the plasmalemma has retracted towards the centre of the cell (Figure 2C), whereas the cell walls of treated melon are strongly and uniformly stained and the plasmalemma is near the cell wall and thus cannot be distinguished in the micrograph (Figure 2D). So, the calcium treatment reinforces and consolidates the cell walls of the melon and preserves cell-to-cell contact at a microstructural level.
Figure 1. Melon stored for 1 week at 4ºC. A, B: Cryo-SEM. C, D: LM. A, C: control melon. B, D: calcium treated melon. is: intercellular space; p: plasmalemma; bcw: broken cell wall.
Figure 2. Melon stored for 3 weeks at 4ºC. A, B: Cryo-SEM. C, D: LM. A, C: control melon. B, D: calcium treated melon. p: plasmalemma. 1. 2. 3.
O. Lamikanra, M.A. Watson, Journal of Food Science 69(6) (2004), pp. 468-472. X. Dong, R.E. Wrolstad, D. Sugar, Journal of Food Science 65(1) (2000), pp. 181-186. The authors are indebted to the Project AGL2003-09208-C03-1 for financial support and to the Universidad Politécnica de Valencia for the grant awarded to L. Alandes.
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Ceramic-loaded mineralizing bioresorbable polymers for orthopaedic applications Linn W. Hobbs1, Tamara Lim1,2, Alexandra Porter3, Hao Wang4, Mark Walton5, and Nicholas J. Cotton6 1. Department of Materials Science & Engineering, Massachusetts Institute of Technology, Cambridge, MA 02139-4307, USA. 2. Department of Materials, Oxford University, Oxford OX1 3PH, UK. 3. Department of Materials, Imperial College London, London SW7 2AZ, UK. 4. Department of Orthopaedic Surgery, Children’s Hospital, Boston, MA 02115, USA. 5. Department of Orthopaedic Surgery, Dunedin School of Medicine, Otago, NZ. 6. Endoscopy Division, Smith & Nephew, Inc., Andover, MA 01810, USA [email protected] Keywords: bioresorbable polymers, mineralization, orthopaedic implants
Bioresorbable polymers are widely used in biomedical applications, for example in surgical sutures and orthopaedic screws. One class of these materials is broadly based on poly(lactide)-poly(glycolide) copolymers, whose proportions govern the rate of physiological resorption. One difficulty encountered in implantation of resorbable polymers into bone is local acidification as the polymer breaks down and the body mounts an inflammatory response. Another is that resorbable polymers are usually neither bioactive nor osteoconductive. This presentation explores a novel bioresorbable polymer-inorganic particulate composite implant that both buffers local acidity and promotes osteoconductive replacement of the implant by bone over a one-year time span. The composite comprises an 85:15 poly-DL-lactide-co-glycolide matrix into which has been incorporated a 35wt% dispersion of calcite (CaCO3) particulates, with a bimodal distribution of micrometer- and sub-micrometer sizes that contribute initial strength (important during insertion) as well as subsequent physiological activity [1,2]. The material was formed into interference screws, used for tibial fixation of replacement tendon in anterior cruciate ligament (ACL) repair [1,2], and into suture anchors for reattaching shoulder ligaments to bone in rotator cuff injuries. The study reported here followed mineralization sequences occurring over a 6-week period in vivo in an ovine animal model undergoing ACL replacement and over a 12week period for model attachment of patella tendon to tibial bone. Progress has been documented using analytical SEM, TEM and STEM of thin sections through the implant and surrounding bone and tendon. Following matrix degradation, the calcite particulates are observed to react with biogenic phosphorus to form calcium phosphate (CP) compounds, with a range of compositions and morphologies, that promote eventual bone infill replacing the initial composite implant (Figure 1). The calcium carbonate buffers local acidity during matrix resorption, not only favoring mineralization but also influencing those CP phases forming.
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Mineralization at tendon-bone interfaces is observed to be more favorably affected by the nearby presence of the calcite particulates, with differences observed in both mineralization morphology and strength of eventual tendon attachment. A mineralized interface layer of what appears to be mineralizing cartilage (Fig. 2), laid down by chrondocytes, forms between and integrates with adjacent bone and tendon. 1.
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W. R. Walsh, N. J. Cotton, P. Stephens, J. E. Brunelle, A. Langdown, J. Auld, F. Vizesi and W. Bruce, Comparison of poly-L-lactide and polylactide carbonate interference screws in an ovine anterior cruciate ligament reconstruction model. J. Arthroscopic and Related Surgery 23 (2007) 757-65. N.J J. Cotton, M. J. Egan and J. E. Brunelle, Composites of poly(DL-lactide-co-glycolide) and calcium carbonate: In vitro evaluation for use in orthopaedic applications. J. Biomed. Mater. Res. Part A (pub;ished on-line 9 August 2007). This work was supported by the Cambridge-MIT Institute, Children’s Hospital Boston, and Smith & Nephew Endoscopy Division.
Figure 1. TEM image of thin section showing bone-like mineralization at 6 weeks in situ in ovine model, following resorption of calcite-loaded bioresorbable polymer interference screw and reaction of calcite particles with biogenic phosphorous. Ca/P ratio is ~1.26. Selected-area electron diffraction pattern (inset) resembles that of hydrodyapatite mineralization in bone.
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Tendon
Figure 2. SEM image of mineralizing cartilagenous layer uniting bone with previously separated tendon after 12 weeks in situ in ovine model in the vicinity of a CaCO3-doped bioresorbable suture anchor. Chondrocyte lacunae are visible in the mineralized layer.
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AFM and TEM study of Ag coated insulin-derived amyloid fibrils M. Gysemans1, J. Snauwaert1, C. Van Haesendonck1, F. Leroux2, B. Goris2, S. Bals2, G. Van Tendeloo2 1. Laboratory of Solid-State Physics and Magnetism, K.U.Leuven, Celestijnenlaan 200D, BE-3001 Leuven, Belgium 2. EMAT, University of Antwerp, Groenenborgerlaan 171, BE-2020 Antwerp, Belgium [email protected] Keywords: nanoparticles, biomaterials
Recently, the use of ordered biological building blocks for the fabrication of nonbiological nanostructures has been explored. Biological molecules offer a variety of new synthetic methods for the development of interesting inorganic materials [1]. It has, e.g., been reported that metallic nanowires are obtained by depositing metals such as Ag or Pd directly onto DNA and tubulin assemblies [2,3]. Here, we rely on the use of insulinderived amyloid fibrils to direct the nucleation, deposition, and assembly of Ag nanoparticles. The fibrils provide an appropriate protein based template for the bottomup fabrication of wires with nanoscale diameter. In order to study deposition patterns of Ag nanoparticles immobilized along insulin-derived amyloid fibrils, we rely on atomic force microscopy (AFM) in combination with transmission electron microscopy (TEM). The formation and deposition of the Ag particles is done by electroless plating of the biomolecular template. The polymerized insulin fibrils are incubated with Ag nitrate, immobilised on a TEM or AFM substrate, and reduced with sodium borohydride. As shown in Figure 1, the AFM images obtained in the “tapping” mode reveal the presence of insulin fibrils with diameters ranging from 3 to 6 nm and dispersedly covered with Ag particles with heights of a few nm. Very few unbound Ag particles can be found in the background (substrate is a piece of an oxidized silicon wafer). The investigated TEM samples are unstained, implying the contrast in the TEM images is dominated by the Ag nanoparticles. Besides agglomerates of Ag particles with micrometer range sizes, chains of nanosize Ag particles are observed. An example of such a Ag nanoparticle chain is shown in Figure 2.a. The chain, indicated by the white arrow, consists of particles with sizes ranging from approximately 10 nm to only a few nm. In this case, the particles do not form a continuous chain. In general, the diameter of the chain corresponds to the diameter of one single particle. The area indicated by a white rectangle in Figure 2.a is shown in more detail in Figure 2.b. For this chain, it seems that the Ag particles are connected to each other and again, the diameter of the chain mainly corresponds to the diameter of one particle. The 3D structure of the Ag coated fibrils can be studied by means of electron tomography, where several projections along different angles serve as an input for 3D reconstruction.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 739–740, DOI: 10.1007/978-3-540-85226-1_370, © Springer-Verlag Berlin Heidelberg 2008
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S. Behrens, W. Habicht, K. Wagner and E. Unger, Adv. Mater. 18 (2006), p.284. K. Keren and E. Braun, Chem. Eng. Technol. 27 (2004), p.447. S. Behrens, K. Rahn, W. Habicht, K.J. Bohm, H. Rosner, E Dinjus and E. Unger, Adv. Mater. 14 (2002), p.1621. This work has been supported by the Interuniversity Attraction Poles (IAP) program (P6/42 on “Quantum effects in clusters and nanowires”).
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Figure 1. AFM images at two different magnifications of insulin fibrils that are dispersedly covered with Ag nanoparticles.
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Figure 2. a) TEM image of a Ag nanoparticle chain. The area indicated by the white rectangle is shown in more detail in b).
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Transmission Electron Microscopy studies of bio-implant interfaces using Focused Ion Beam microscopy for sample preparation F. Lindberg1, A. Palmquist2, L. Emanuelsson2, J. Heinrichs1, R. Brånemark3, F. Ericson1, P. Thomsen2 and H. Engqvist1 1. Materials Science, Department of Engineering Sciences, Ångström Laboratory Uppsala University, Uppsala, Sweden. 2. Department of Biomaterials, Sahlgrenska Academy at the University of Gothenburg, Göteborg, Sweden. 3. Department of Orthopaedics, Sahlgrenska University Hospital, Göteborg, Sweden. [email protected] Keywords: TEM, FIB, biocompatibility, titania surfaces, implants, bio-implant interfaces.
Biocompatibility of metallic implants is governed by surface characteristics. TEM sample preparation using FIB microscopy very much facilitates our research concerning implant-tissue interfaces as well as studies of bioactive surfaces. This talk will cover some of our research on bioactive titanium surfaces and interfaces of titanium and bone. Hydroxyapatite apatite precipitation on single crystalline rutile structured TiO2[1]. Hydroxyapatite precipitates faster structured TiO2 (rutile or anatase) compared to amorphous [2]. Important in this context is the ordered arrangement of hydroxy groups that are attached to the titania surface in solution. This calls for epitaxial growth of hydroxyapatite on ordered TiO2. However, such growth has to date not been observed. In this work we grew hydroxyapatite on different crystalline faces of rutile from simulated body fluid. With X-ray diffraction it could be concluded that hydroxyapatite precipitated faster on a (001) crystal surface compared to (110) and (100). From HRTEM images obtained from FIB prepared sample slices it could be concluded that: • Hydroxyapatite precipitates at the immediate surface. • FFT analysis of the micrographs reveals that the hydroxyapatite orients differently on different rutile crystal faces, see figure 1. Considering the growth direction of the crystallites the speed of growth could be explained. Nanostructuring and preparation of titania surfaces using laser [3]. Figure 2a shows a titanium surface that has been subjected to laser treatment. In addition to the formation of solidified droplets due to melting and rapid cooling, a nanostructured titanium oxide surface film is created. Laser modified implants demonstrate a significantly increased stronger bone anchorage compared to turned machined implants. From TEM investigations of the bio-implant interfaces it could be concluded that laser treatment of titanium produces a direct contact between the implant surface and mineralized bone with bone growing inside the nanostructured region, see figures 2b. 1.
Lindberg F, Heinrichs J, Ericson F, Thomsen P, Engqvist H, Submitted 2008.
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Uchida M, Kim HM, Kokubo T, Fujibashi S, Nakamura T, J Biomed Mater Res A 2003;64A:164-170. Palmquist A, Lindberg F, Brånemark R, Engqvist H, Thomsen P. Submitted 2008.
Figure 1. HRTEM micrographs (left part) and the corresponding FFTs of hydroxyapatite precipitated on the (001) and (110) rutile surfaces. The arcs of points representing the 100 hydroxyapatite planes indicate textural growth, different on the two surfaces.
Figure 2. a) Laser treatment of a Ti surface creates macrostructured surface due to melting and rapid cooling but also a nanostructured surface oxide on which bone growth is facilitated (b).
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AFM and SEM of Wax Crystallisation on Artificial Surfaces Controlled by Temperature and Solvents A. Niemietz1, W. Barthlott1, K. Wandelt2 and K. Koch3 1. Nees Institute for Biodiversity of Plants, Rheinische Friedrich-Wilhelms-Universität Bonn, University Meckenheimer Allee 170, 53115 Bonn, Germany 2. Institute for Physical and Theoretical Chemistry, Rheinische Friedrich-WilhelmsUniversität Bonn, Wegelerstr. 12, 53115 Bonn, Germany 3. Nanotribology Laboratory for Information Storage and MEMS/NEMS, Ohio State University, 201 West 19th Avenue, Columbus, Ohio 43210-1142, USA [email protected] Keywords: wax crystallisation, self assembly, structured surfaces
Plant surfaces show a great variety of functional structures between micro- to nanodimensions. Epicuticular wax crystals are very frequent and cause extreme water repellence and are responsible for the self-cleaning property. These waxes are complex mixtures of aliphatic components like primary or secondary alcohols, aldehydes and fatty acids, which self assemble into three-dimensional structures on the plant surface as well as on artificial surfaces. We used atomic force microscopy (AFM) and scanning electron microscopy (SEM) to study the process of wax crystal formation under varying environmental conditions (temperature and solvent) to optimize the development of biomimetic hydrophobic structured surfaces. To this end different wax crystal morphologies e.g. tubules and platelets [1], have been separated from the plant surface by dissolving them in chloroform, followed by chemical analysis and recrystallisation from the solution. Earlier experiments showed that recrystallisation out of solutions leads to unequal mass distributions, with typical ring patterns (coffee-drop-effect) on the substrates. Therefore the waxes were applied by physical vapour deposition (PVD). This method led to a homogenous distribution of the wax structures on the substrates and perfect crystals of octacosan-1-ol wax platelets, were grown at standard conditions. Structures which are composed of more than one compound, e.g. the nonacosan-10-ol wax tubules from the Lotus (Nelumbo nucifera) [3], needed a special treatment for recrystallisation. After application of tubules waxes by PVD a layer of amorphous structures without any reassembly was formed [Fig. 1a]. Increasing temperature and/or using solvent vapours led to an increased mobility of the wax molecules resulting in a homogenous distribution of wax tubules along the surface [Fig. 1b]. In general it is to conclude that physical vapour deposition is a suitable method for the generation of wax morphologies comparable to those know from plant surfaces. AFM analysis of the crystallisation process provides information about the interactions between the wax-molecules- and the substrates and the wax-molecules and solvents used for their mobilisation. In future these biomimetic wax surfaces could be used for
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e.g. high resolution investigations of liquid interactions at waxy surfaces, as for example the spreading of liquids on hydrophobic surfaces. 1 2 3
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Barthlott W, Neinhuis C, Cutler D, Ditsch F, Meusel I, Theisen I, Wilhelmi H (1998) Classification and terminology of plant epicuticular waxes. J. Linn. Soc. 126, 137-260. Koch K, Neinhuis C, Ensikat HJ, Barthlott W (2004) Self assembly of epicuticular waxes on plant surfaces investigated by Atomic Force Microscopy (AFM). J Exp Bot 55: 711-718. DOI:10.1093/jxb/erh077 Barthlott W, Neinhuis C (1997) Purity of the sacred lotus or escape from contamination in biological surfaces. Planta 202, 1-7.
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Figure 1: SEM pictures of wax on glass after application by PVD a) without changing environmental conditions b) reassembled wax tubules after increasing temperature
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Characterization of layer-by-layer microcapsules made of hyaluronic acid by CLSM, SEM and TEM I. Pignot-Paintrand1, A. Szarpak1,2 and R. Auzely-Velty1 1. Centre de Recherches sur les Macromolécules Végétales (CERMAV-CNRS), BP53, 38041 Grenoble cedex 9, France, affiliated with Université Joseph Fourier and member of the Institut de Chimie Moléculaire de Grenoble 2. Jan Dlugosz University, Armii Krajowej 13/15 Ave., Czestochowa, Poland [email protected] Keywords: hyaluronic acid, microcapsules, SEM, TEM
During the past decades a large variety of micro- and nanocarriers have been developed in order to improve efficiency, availability and toxicity profiles of drugs. Hollow capsules prepared by the layer-by-layer (LbL) assembly of oppositely charged polyelectrolytes on a colloidal template, followed by its decomposition, have recently emerged as attractive vehicles in the field of drug delivery. We chose to design microcapsules from hyaluronic acid (HA) and poly(allylamine hydrochloride) (PAH) using CaCO3 microtemplates. Hyaluronic acid (HA) which is a natural linear polysaccharide belonging to the glycosaminoglycan family, exhibits excellent biocompatibility and biodegradability which can be advantageously exploited for many drug delivery applications [1]. Polyelectrolyte multilayers based on biopolymers are hydrogels and must be considered as "soft" and sensitive materials [2]. Having HA samples with different molecular weights in hands, we investigated their effect on the formation of polyelectrolyte films. To characterize colloidal templates, we used scanning electron microscopy (SEM) The porous structure of spherical microparticles is shown in Figure 1. The surface coating was observed by scanning electron microscopy and transmission electron microscopy (TEM). For transmission electron microscopy, the particles coated with two to six layers of HA/PAH were embedded in Lowicryl HM20 polymerized at 22°C under indirect UV light, before sectioning. The use of HA of different molecular weights and at various concentration during multilayer formation influences the thickness of the film (Figure 2). After the adsorption of bilayers with the desired number was achieved, the calcium carbonate core was subjected to decomposition. In order to gain insight into the composition of the multilayer walls, we examined the capsules by confocal laserscanning fluorescence microscopy (CLSM) after the labelling polymers (Figure 3). 1. 2. 3.
N.E. Larsen, E.A. Balazs, Adv. Drug Delivery Rev. 7 (1991), p. 279. A. Engler, L. Richert, J. Wong, C. Picard, D. Discher, Surface Science 570, (2004), p. 142. We kindly acknowledge Frédéric Charlot at CMTC-INPG, Grenoble, France, for his help with SEM observations using the Zeiss ultra 55 FEG-SEM
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Figure 1. SEM image of colloidal template. Insert: high-resolution SEM image of porous structure of spherical CaCO3 microparticles (FEG-SEM, in lens detector).
Figure 2. TEM images of two different concentration of HA11, a) 5g/L and b) 14g/L.
Figure 3. CLSM images of hollow capsules
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Atomic force microscopy analysis of crystalline silicon functionalization with oligonucleotides A. Ponzoni1, G. Faglia1, M. Ferroni1, G. Sberveglieri1, A. Flamini2, G. Andreano3, L. Cellai3 1. CNR-INFM, Dipartimento di Chimica e Fisica per l'Ingegneria e per i Materiali Universita' di Brescia , Via Valotti 9, Brescia , Italy 2. Istituto di Struttura della Materia and 3. Istituto di Cristallografia, CNR, Monterotondo Stazione, Rome, Italy [email protected] Keywords: atomic force microscopy, oligonucleotide, silicon
The analysis of oligonucleotide (ODN) loading over inorganic substrates is of fundamental importance to properly develop organic-inorganic systems such as biosensors. For this purpose, silicon is one of the most suitable substrates, but few works have been reported compared to other substrates such, for example, oxides, glasses or gold [1]. Crystalline silicon functionalization is a delicate process since the use of both basic and aggressive acidic solutions can damage the Si surface. A strict control of the functionalization process is of crucial importance both to preserve the high quality of the substrate surface (and its properties) and to optimize the final device performance. Functional parameters such as sensitivity and selectivity of the biosensor, depends on the ODN probe-target hybridization efficiency, which in turn depends on the optimization of ODN probe loading on the Si substrate. In this work we focus on the development of a simple methodology for the immobilization of controlled amounts of oligonucleotides on unoxidized crystalline Si(100) surface through a stepwise procedure and using a compact monolayer of 6heptynoic acid, as a spacer molecule. To monitor this procedure, we make use of fluorescence and atomic force microscopy (AFM) techniques. In particular we use fluorescence as the standard method to characterize loading and hybridization from a macroscopic point of view as previously described [1] and AFM techniques, namely topography and Phase Detection Microscopy (PDM), to characterize the functionalized substrates at nanometric level obtaining direct imaging of anchored ODNs. The ODN loading process consists in the following steps. At first freshly hydrogenated silicon is put in the contact with neat 6-heptynoic acid ethyl ester, under nitrogen, in a dry-box, for two days, at room temperature; as a result, the silicon surface becomes covered by a monolayer of the ester, covalently anchored to it through the C-Si bond. Then the ester-terminated surface was converted to the corresponding carboxylic acid via acidic hydrolysis in 3 M HCl, for two days, at room temperature, in air. Finally amino-terminated ODNs are immobilized onto this surface through an amidation in aqueous solution, upon incubation, at room temperature, for a variable length of time (30-300 min). In this way, several silicon samples containing ODNs at variable density on their surface were obtained.
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Topographic AFM and Phase Detection Microscopy (PDM) have been used to visualize the biological samples over the silicon substrate, prepared as described above. Morphological analysis of samples with low ODN density reveal a low substrate roughness confirming that the proposed functionalization approach does not damage the Si substrate. Isolated ODNs vertically standing on the substrate are imaged as particles with height of about 4 nm and full with at half height (FWHM) of about 12 nm (Figure 1, a). In AFM (tapping mode) measurements, the particle-tip convolution and the compressibility of ODNs cause height underestimation and width is overestimation. Furthermore, it has been reported that ODNs can bind to the substrate forming different angles and thus exhibiting reduced height and increased FWHM. These phenomena may induce errors in ODN identification by means of topography analysis alone. So far, PDM, as a technique sensitive to the material properties and suitable to discriminate between soft matter (ODNs) and the hard substrate (Si), has been employed to enhance the capability to identify ODNs. An example is reported in Figure 1, b. ODNs are distinguished over the Si substrate thank to the different elastic properties and are imaged as dark areas (lower phase leg values). In some cases, PDM signal provides a better contrast then the topography one. The combination of techniques used (AFM, PDM and fluorescence) allows to get a reliable characterization of the proposed functionalization method together with visualization of ODNs displacements over Si (100). 1. 2.
F. Cattaruzza et al. Nucleic Acids Research, 34, 2006, p.1-13 This work has been partly supported by the Ministero dell'Istruzione, dell'Università e della Ricerca PRIN 2005 "Quasi mono dimensional nanosensors for label free ultra sensitive biological detection". Prof. Diligenti (University of Pisa, Phys. Dept.) is gratefully acknowledged for Si (100) substrates.
Figure 1. AFM 3D image of an isolated couple of ODNs (a) and 2D phase image (b) of a sample featuring a high ODN density, identified by dark color (low phase values).
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Hidden hierarchy of microfibrils within fluorapatite gelatine nanocomposites induced by intrinsic electric dipole fields Paul Simon1 and Rüdiger Kniep1 1. Max Planck Institute for Chemical Physics of Solids, Nöthnitzer Str. 40, 01187 Dresden, Germany [email protected] Keywords: biomineralisation, electric fields, apatite-gelatine, bio-composites
The biomimetic system fluorapatite-gelatine bears strong resemblance to the biosystem hydroxyapatite-collagen which plays a decisive role in the human body as functional material in the form of bone and teeth. Initial stages of growth comprise micrometer-sized, hexagonal prismatic fluorapatite-gelatine composite seeds (marked violet in Figure 1a). These species develop in subsequent growth stages to dumbbells states and complete their development as closed, notched spheres. We focused our efforts on the central composite seed (Figure 1b and Figure 2, top), since it represents the initial and thus the fundamental growth step during the morphogenesis [1]. Longitudinal focused ion beam thin cuts ([ ׀׀001] of apatite) of young seeds with a perfect hexagonal-prismatic habit were prepared (Figure 2, top right). The zoomed TEM image (Figure 2, bottom left) clearly shows that the 3D nanocomposite arrangement is dramatically over lied by a pattern consisting of gelatine micro-fibrils with diameters scaling around 10 nm. The variations in orientation and concentration of the microfibrils lead to a spatial subdivision of the young seed consisting of three distinct areas. (1) Cone-like area with a low concentration of microfibrils running perpendicular to the basal plane. (2) Area with bent microfibrils and extending in direction of edges between basal and prism faces. (3) Area with the largest parts of the prism faces forming its outer boundaries. Micro-fibrils arrange themselves to finally find an orientation perpendicular to the prism faces. We assume that gelatine triple-helices exhibit opposite charges at their ends. By adding up all these microscopic dipoles a macroscopic electric dipole of the composite is formed as visualized by electron holography (Figure 1c) [2]. This could have a decisive influence on further growth steps and thus on the morphogenesis. The intrinsic pattern of gelatine microfibrils embedded within the periodic matrix of the fluorapatite-gelatine-nanocomposite represents an increased level of complexity providing the conception for future form-developments and being present already within a young composite seed with its perfect hexagonal-prismatic habit (Figure 2, bottom right) [3]. 1. 2. 3.
R. Kniep, S. Busch, Angew. Chem., 108 (1996), 2787-2791. Angew. Chem. Int. Ed Engl. 35 (1996), 2624-2626. P. Simon, D. Zahn, H. Lichte, R. Kniep, Angew. Chem. Int. Ed. 45 (2006), 1911-1915. R. Kniep, P. Simon, Angew. Chem. Int. Ed. 47 (2008), 1405-1409.
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Figure 1. (a) Superposition of growth stages of fluorapatit-gelatine nanocomposite aggregates (SEM images). (b) Hexagonal prismatic nanocomposite seed; conventional TEM micrograph. (c) Retrieved phase image of an electron hologram exhibits the electric potential distribution around a seed. The observed projected potential corresponds to a mesoscopic dipole.
Figure 2. (left top) Scanning ion image of a virgin composite seed before FIB thinning is showing the perfect hexagonal symmetry of a fluorapatite single crystal. (right top) Overview TEM micrograph of a longitudinal FIB thin cut of the young seed. (bottom, left) Enlarged section revealing three distinct areas within the young seed (red lines are suggesting the borders between the different regions). (bottom, right) Bunch of bent micro-fibrils (red/dotted lines) within a “young” seed (black outline) and extending to the “mature” seed (blue outline).
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Self-assembled block copolymer structures studied by transmission electron microtomography H. Jinnai1, T. Kaneko1, C. Abetz2, V. Abetz2 1. Department of Macromolecular Science and Engineering, Graduate School of Science and Engineering, Kyoto Institute of Technology, Kyoto Institute of Technology, Kyoto 606-8585, Japan 2. Institute of Polymer Research, GKSS Research Centre Geesthacht GmbH, 21502 Geesthacht, Germany [email protected] Keywords: block copolymers, transmission electron tomography, double helical structure
Block copolymers, consisting of multiple, chemically dissimilar sequences covalently linked together, exhibit highly periodic nanoscale structures [1]. The nanoscale structures of the diblock copolymers include spheres of A(B) on a bodycentered cubic lattice in a B(A) matrix, cylinders of A(B) on a hexagonal lattice in a B(A) matrix, several complexes (bicontinuous), called gyroid, of A(B) in a B(A) matrix, and co-alternating lamellae. With the rapid advancement in high-precision synthesis, the nanostructures have become relatively easily tailored by molecular engineering, which is very appealing for the applications of nanotechnology [2], e.g., low-k films, nanoporous membranes, and high-density information storage media. In some cases, due to their sophisticated self-assembling capabilities, the supramolecules were used to mimic a well-known biological architecture, the helical structure. One of the most fascinating and fundamental helical structures in nature is formed by DNA, a right-handed double helical structure, while an important motif in proteins is another important helix (α-helix), which is a right-handed single helix. Collagen, as another significant example, forms a triplehelix. Cornelissen et al. synthesized an amphiphilic block copolymer containing polystyrene and a charged helical polyisocyanide head-group derived from isocyano-L-alanine-L-alanine and isocyano-L-alanine-L-histidine. They observed the helical structure in a solution with sodium acetate [3]. In 2004, Ho et al. demonstrated that their polystyrene-block-poly(Llactide) formed hexagonally-packed poly(L-lactiede) nanohelices with a left-handed helical structure in the bulk, which is a more suitable state for nanotechnology applications than the solution state [4]. In the above examples, the chiral entity of the constituent blocks plays an important role in the formation of the helical superstructures. Thus, the chirality of the chemical compounds remains to be one of the main driving forces for the formation of the helical structure. An interesting nanohelical structure, helical coils around a cylinder, has been observed in ternary block copolymers in the bulk [5,6]. The chiral superstructures were realized by the relative incompatibilities and the block chain conformation inside the structure, not by the chirality on a segmental scale. However, all of the previous attempts failed to claim a “double” helical structure.
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In the present study, we show three-dimensional (3D) structure of the “achiral” triblock terpolymer, poly(styrene-block-butadiene- block -(methyl methacrylate)), in the bulk using high-resolution transmission electron microtomography (TEMT) [7,8]. The “double” helical structure, hexagonally-packed polystyrene (PS) cylinders with polybutadiene (PB) helices around them in the poly(methyl methacrylate) (PMMA) matrix, were formed. The fraction of PB microdomains with a left-handed helical sense to right-handed one was similar, but the opposite helical sense between adjacent double helical structures seemed to be energetically favoured. This phenomenon was discussed in terms of the packing frustration of the block chains inside the nanoscale structure. [9] 1. 2. 3. 4. 5. 6. 7. 8. 9.
E. L. Thomas, D. M. Anderson, C. S. Henkee and D. Hoffman, Nature 334 (1998), p. 598. I. W. Hamley, Nanotechnology 14 (2003), p. R39. J. Cornelissen, M. Fischer, N. Sommerdijk and R. J. Norte, Science 280 (1998), p. 1472. R. M. Ho et al., J. Am. Chem. Soc. 126 (2004), p. 2704. U. Krappe, R. Stadler and I. Voigt-Martin, Macromolecules 28 (1995), p. 4558. U. Breiner, U. Krappe, V. Abetz and R. Stadler, Macromol. Chem. Phys. 198 (1997), p. 1051. H. Jinnai et al., Phys. Rev. Lett. 84 (2000), p. 518. H. Jinnai, Y. Nishikawa, T. Ikehara and T. Nishi, Adv. Polym. Sci. 170 (2004), p. 115. The authors are grateful to NEDO for support through the Japanese National Project “ Nano-Structured Polymer Project” by the Ministry of Economy, Trade and Industry and for support from the Ministry of Education, Science, Sports and Culture through Grants-in-Aid No. 1855019 and No. 19031016.
Figure 1. TEM images of the SBM triblock terpolymer. (a) Top and (b) side views of the PS cylinders. The pattern of the SBM triblock terpolymer is schematically drawn in the insets. OsO4-stained PB microdomains are shown. Bar indicates 200 nm. (c) 3D reconstructed image of a left-handed double helical structure marked by the dashed rectangular region in part (b). Blue and red domains correspond to the PB microdomains.
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Quantitative chemistry and orientation of polymers in 2-d and 3-d by scanning transmission X-ray microscopy A.P. Hitchcock1, G.A. Johansson1, D. Hernández Cruz1, E. Najafi1, J. Li1 and H. Stöver2 1.
Brockhouse Institute for Materials Research, McMaster University, Hamilton, ON, Canada L8S 4M1
[email protected] Keywords: polymers, nanotubes, STXM, NEXAFS, X-ray absorption, tomography
In the last 20 years, synchrotron based soft X-ray scanning transmission X-ray microscopy (STXM) [1,2] has emerged as a powerful micro-analytical technique with particular strengths in the analysis of soft matter – polymers, biological and environmental samples. Its analytical power stems from spatially resolved near-edge Xray absorption spectroscopy (NEXAFS); thus it produces information analogous to that delivered by electron energy loss spectroscopy in analytical transmission electron microscopes (TEM-EELS). NEXAFS microscopy has significant advantages for studies of soft matter, which is typically a challenge for TEM-EELS due to radiation damage. This presentation will describe current state-of-the-art soft X-ray microscopes (including the recently commissioned spectromicroscopy facility at the Canadian Light Source in Saskatoon), and illustrate their performance with examples of twodimensional and three-dimensional quantitative studies of chemical and orientational information in a variety of systems, including partially and fully hydrated samples. STXM is a microprobe technique which uses a Fresnel zone plate (ZP) to focus monochromated synchrotron X-rays to a point with a diameter of 20-30 nm and a depth of focus of 500-1500 nm (depending on photon energy and details of the ZP). In most current implementations the sample is raster scanned through the focal point while detecting the X-rays transmitted through the sample. The photon energy is scanned and spectral information is obtained though point, line, or image sequences. The focal length varies linearly with photon energy so precise control of ZP to sample distance and ZPsample alignment is needed. This is achieved optimally with interferometric control [3]. Polymers, particularly those with aliphatic regions and carbonyl functional groups, are challenging to study by TEM-EELS due to radiation damage. We (and many other groups) are using STXM to investigate a wide range of synthetic and natural polymer systems, including controlled release polyurea capsules [4], polyurethane foams, coreshell microspheres, etc. In addition to chemical mapping, the linearly polarized synchrotron radiation allows spatially resolved measurements of dichroism, either by moving the sample [5] or by using an elliptically polarizing undulator (EPU), which can rotate the plane of polarization over a full 180o. The latter approach has proved particularly powerful in analysis of the spatial distributions of β-sheet orientation in silk fibers [6]. Recently we have implemented a controlled humidity cell which is being used to investigate the impact of humidification on the orientational ordering of silk fibers. EPU polarization studies are also being used to probe the quality of multi-walled
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carbon nanotubes (MWCNTs). The C 1s → π* transition (285 eV) of CNTs has a very strong dichroic response. In a perfect CNT, that signal is zero when the E-vector is parallel to the tube axis and maximum when the E-vector is perpendicular to the tube axis (see Figure 1). When MWCNTs of differing structural quality are compared significant differences in the π* dichroic response are observed which can be related to the local level of sp2 defects. Within STXM spatial resolution limits, this may provide a useful tool to evaluate and assist optimization of electronic devices based on individual CNTs. In addition to 2-d projection analysis, it is often important to examine chemical structure in 3-dimensions. 3-d STXM has been achieved with serial sectioning in toner particles [8] and, with angle-scan computed tomography, in frozen tissue [9], and in wet biofilms and wet latex microspheres [10]. Figure 2 displays a rendering of 3-d quantitative distributions of a very dilute polyacrylate gel (~2% dry weight), inside submicron diameter polystyrene microspheres in water inside a glass capillary. The penetrating capability of the sub-O 1s edge X-rays was critical to achieve this result. 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11.
A.P. Hitchcock et al, J. El. Spect. Rel. Phen. 144-147 (2005) p. 259. H. Ade and A.P. Hitchcock, Polymer 49 (2008) p. 643. A.L.D. Kilcoyne et al., J. Synchrotron Rad. 10 (2003) p. 125. A.P. Hitchcock, et al., J. El. Spect. Rel. Phen. 156-158 (2007) p. 467. D. Hernández Cruz et al., Rev. Sci. Inst. 78 (2007) p. 033703. M.E. Rousseau et al. J. Am. Chem. Soc. 129 (2007) p. 3897. E. Najafi et al, Small (2008) submitted. A.P. Hitchcock et al., J. Phys. IV France 104 (2003) p. 509. Y. Wang, C. Jacobsen, J. Maser and A. Osanna, J. Microscopy, 197 (2000) p. 80. G.A. Johansson et al. , J. Synchrotron Rad. 14 (2007) p. 395. Research supported by NSERC, CFI and Canada Research Chair. STXM carried out at ALS (supported by U.S. Dept. of Energy) and CLS (supported by NSERC, CIHR, UoS).
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Figure 1. (upper) image (285 eV) and dichroism in indicated area of high quality arc discharge multiwalled carbon nanotube (AD-MWCNT). (lower) image (285 eV) and dichroism in indicated area of a lower quality chemical vapour deposition sample (CVD- MWCNT).
Figure 2. (upper left) NEXAFS basis for mapping polyacrylate in polystyrene latex microspheres in water. (right) rendering of quantitative 3-d distributions of polyacrylate derived from the difference of tomograms (60 angles from 0 – 180o) measured at 530 and 532 eV. The PS spheres are 800 nm in diameter.
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Characterization of cavitation processes in filled semi-crystalline polymers F. Addiego1, J. Di Martino1, D. Ruch1, A. Dahoun2 and O. Godard2 1 - Laboratoire de Technologies Industrielles, Centre de Recherche Public Henri Tudor, 66 Rue de Luxembourg, BP 144, L-4002 Esch-sur-Alzette, Luxembourg 2 – Laboratoire de Physique des Matériaux, Institut National Polytechnique de Lorraine – Nancy Université, Parc de Saurupt, 54042 Nancy Cedex, France [email protected] Keywords: polyethylene, composites, cavitation
Modification with rigid filler particles of semi-crystalline polymer has received considerable attention in recent years because it is an easy and cheap method to enhance impact toughness of pure matrix at low temperature. It is generally admitted that improvement of toughness is linked to the formation of voids by matrix/particle interface debonding, which facilitates molecular rearrangement of interparticle ligaments under stress [1]. But, characterization of these debonding mechanisms is qualitative to date. This is firstly due to the difficulties encountered to estimate volume change induced by void development. Indeed, quantification of volume variation during mechanical tests was a challenge from necking initiation that causes a localization of mechanical variables. Thanks to recent progress in optical extensometers (VidéoTraction [2]), it is now possible to record volume strain up to large deformation. Secondly, effect of sample preparation on scanning electron microscope (SEM) observation has been neglected. Attention has systematically been focused on one preparation mode (for example: cryofractured sample or polished specimen + chemical etching [3], followed or not by metal coating). Morphological and/or chemical artifacts may arise from preparation mode and consequently influence the interpretation. The aims of this paper is to study the influence of particle size (micro and nano) on volume variation and to characterize cavitation mechanisms from various methods including polishing, cryofracture, cryosurfacing and fine lamella preparation (for ESEM transmission mode) followed or not by a chemical etching. We analyze cavitation processes of high-density polyethylene (HDPE) filled with 5w% of micro (1 µm) and nano (40-70 nm) calcium carbonate (CaCO3). The materials are subjected to tensile tests at ambient temperature and under constant strain rate by the means of VidéoTraction device, whereas post-mortem investigations are carried out with an environmental scanning electron microscope ESEM FEI Quanta 200. The first results indicate that fillers and their size have an important influence on volume strain of HDPE (figure 1, note that a compaction effect is observed in the nanocomposite, this effect is probably due to an overall underestimation of volume strain). Cryosurfaced morphology of the nanocomposite shows that debonding mechanisms are active in this material (figure 2). Complementary investigation methods are currently in progress for each material.
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Figure 1. True stress – true strain and volume strain – true strain curves (20 °C, 1.103 s-1) of HDPE, HDPE + 5 w% CaCO3 micro, HDPE + 5 w% CaCO3 nano (materials are loading up to a strain 1.0 and subsequently unloading)
Figure 2. Cryosurfaced HDPE + 5 w% nano CaCO3 (residual axial strain nearly 0.8) 1 2 3
A. Lazzeri, Y.S. Thio and R.E. Cohen, Journal of Applied Polymer Science 91 (2002), p. 925 Y.L. Yang, S.L. Bai, C. G’Sell and J.M. Hiver, Polymer Engineering and Science 46 (2006), p. 1512 F. Addiego, A. Dahoun, C. G’Sell and J.M. Hiver, Polymer 47 (2006), p. 4387
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Quantitative analysis of protein gel structure by confocal laser scanning microscopy Komla Ako, Lydiane Bécu, Taco Nicolai, Jean-Cristophe Gimel, Dominique Durand Polymères, Colloïdes, Interfaces, UMR CNRS, Université du Maine, 72085 Le Mans Cedex 9, France [email protected] Keywords: Quantitative, gel structure, confocal
Scattering techniques are suited to quantitatively study the structure of protein gels [1, 2] at length-scales smaller than one micron, but it is clear from microscopy images that protein gels may exhibit structural features at larger length scales [3]. Microscopy has been used for qualitative comparisons of protein gels formed at different conditions, but little has been done to extract quantitative information comparable to that yielded by scattering methods. We have developed a novel approach to study the microscopic structure of protein gels and aggregates by quantitative analysis of images obtained by confocal laser scanning microscopy (CLSM). The pair correlation function and the structure factor of the images were calculated numerically "Figure 1." We have established that if proper care is taken the pixel intensities of the images are proportional to the protein concentration. Therefore the result of this analysis is equivalent to scattering measurements, but on larger length scales "Figure 2." We will show that for isotropic systems, such as protein gels, the analysis of 2-D and 3-D images give identical results, but the spatial resolution of the former is better. We used confocal microscopy to quantify the effect of electrostatic interaction on the structure of protein gels by varying the pH and the ionic strength. This study extends a recent investigation using scattering techniques [1, 4, 5] to more heterogeneous gels formed closer to the isoelectric point or in the presence of more salt. We also studied mixed protein/polysaccharide gels for which the structure is determined by the competition between aggregation and phase separation [6, 7]. 1. 2. 3. 4. 5. 6. 7.
Durand, D.; Gimel, J.C.; Nicolai, T. (2002), Physica A, 304, 253. Nicolai, T.; Durand, D. (2007) Curr. Opin. Colloid Interface Sci., 12, 23 Olsson, C.; Langton, M.; Hermansson, A.-M. (2002) Food Hydrocolloids, 16, 111 Nicolai T. In Food Colloids, Self-Assembly and Material Science; Dickinson, E.; Leser, M. E., Eds.; RSC Publishing: Cambridge, 2007; pp 35-56. Mehalebi, S.; Nicolai, T.; Durand, D. to be published Baussay, K.; Nicolai, T.; Durand, D. (2006) Biomacromolecules, 7, 304. Baussay, K.; Nicolai, T.; Durand, D. (2006) J. Coll. Int. Sci, 304, 335
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Figure 1. (Left) Semi-logarithmic representation of the pair correlation functions of heat-set β-lg gels formed at pH 7 and different NaCl concentrations. The solid lines β
represent fits to g ( r ) = B1 exp[−( r / ξ ) ] + 1 . (Right) Double logarithmic representation of the master curve of the data shown at the left hand obtained by normalizing (g(r)-1)with the amplitude and r with the correlation length. The straight line gives an approximate description of the data over a limited range of r and could mistakenly be interpreted as a self similar structure. The inset shows (g(r) -1)/B1.
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Figure 2. Dependence of the correlation length of heat-set β-lg gels formed at pH 7 on the NaCl concentration. Circles represent data obtained with light scattering and squares represent data obtained with CSLM.
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Thermal stability of organic solar cells: the decay in photocurrent linked with changes in active layer morphology S. Bertho1, I. Haeldermans1, A. Swinnen1, J. D’Haen1, L. Lutsen2, T.J. Cleij1, J. Manca1,2 and D. Vanderzande1,2 1. Hasselt University, Institute for Materials Research, Wetenschapspark 1, Diepenbeek, B-3590, Belgium 2. IMECvzw, IMOMEC, Diepenbeek, B-3590, Belgium [email protected] Keywords: organic solar cells, degradation, TEM
Nowadays bulk heterojunction polymer:fullerene (PCBM) solar cells reach efficiencies of 5% through the use of high mobility donor polymers (e.g. P3HT) and through a continued nanoscale control of the morphology of the donor-acceptor interpenetrating networks [1]. One of the general bottlenecks of organic solar cells is their poor stability. Organic solar cells have a low resistance towards oxygen, UV-light, high temperatures etc. This work focuses on the thermal stability of organic solar cells. We compare the thermal stability of 3 different polymer:PCBM bulk heterojunction solar cells [2,3]. The polymers under investigation are poly [ 2 - methoxy - 5 - (3’ , 7’ – dimethyloctyloxy) -1,4-phenylene vinylene] (MDMO-PPV) with a glass transition temperature (Tg) of about 50°C, poly(3-hexylthiopene) (P3HT) with a Tg of about 6°C and ‘High Tg PPV’, a polymer designed by Merck OLED with a Tg of about 138°C. The effect of a thermal treatment on the photocurrent output of the solar cells, shows a clear difference in degradation behavior for the different materials under investigation (Figure 1). The photocurrent output of solar cells based on MDMO-PPV shows a gradual, temperature dependent decay. The decay characteristics of solar cells based on P3HT are completely different; for short-time annealing, the efficiency increases but for longtime annealing the decay also decreases. For solar cells based on ‘High Tg PPV’, a rather stable behavior is observed, without significant impact of thermal treatment. The effect of a thermal treatment on the photocurrent output is reflected in the active layer morphology, which is studied with Transmission Electron Microscopy. The active layers based on MDMO-PPV show a rapid formation of PCBM-clusters upon annealing (Figure 2). The lower interfacial area between electron acceptor and electron donor results in a lower photocurrent output. P3HT shows a dual crystallization behaviour (Figure 3): on one hand, P3HT crystallizes, which leads to improved charge conduction and a higher photocurrent output; on the other hand, PCBM groups into clusters, again reducing the interfacial area between electron acceptor and electron donor, which results in the photocurrent decrease. Within the active layers based on ‘High Tg PPV’ hardly any morphology changes are observed upon annealing (Figure2). This explains the better thermal stability of the photocurrent output of this material.
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1. 2. 3. 4.
W. Ma, C. Yang, X. Gong, K. Lee, A.J. Heeger, Adv. Funct. Mater. 2005, 15, 1617 S. Bertho, I. Haeldermans, A. Swinnen, W. Moons, T. Martens, L. Lutsen, D. Vanderzande, J. Manca, A. Senes, A. Bonfiglio, Sol. En. Mater. Sol. Cells 2007, 91, 385. S. Bertho, G. Janssen, T. J. Cleij, B. Conings, W. Moons, A. Gadisa, J. D’Haen, E. Govaerts, L. Lutsen, J. Manca, D. Vanderzande, in press at Sol. En. Mater. Sol. Cells The research was carried out in the framework of the IWT-project 030220 “Nanosolar”, the FWO-project G.0252.04 and the interregional project OLED+. The work, as part of the project “SOHYDs” within the European Science Foundation EUROCORES Programme was also supported from funds by the FWO (G.0685.06) and the EC Sixth Framework Programme, under contract N. ERAS-CT-2003-980409. S. Bertho is research assistant of the Fund for Scientific Research, Flanders (Belgium) (F.W.O.). We thank Dr. H. Becker of Merck OLED Materials GmbH for the supply of the ‘High Tg PPV’.
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Figure 1. Relative decay of the photocurrent at an annealing temperature of 110°C for solar cells based on MDMOPPV:PCBM (1:4), P3HT:PCBM (1:1) and ‘High Tg PPV’:PCBM (1:4).
Figure 2. TEM images of MDMOPPV:PCBM 1:4 blends (a-b) and ‘High Tg PPV’:PCBM 1:4 blends (a’-b’). The blends were annealed at 110°C for 0h (a, a’), 16h (b, b’), yielding formation of large PCBM-clusters for the MDMOPPV:PCBM blends while maintaining a more stable morphology for the ‘High Tg PPV’:PCBM blends (scale bar: 2μm)
Figure 3. TEM images of a P3HT:PCBM 1:1 blend annealed for 16h at 110°C. PCBMclusters (left) are formed in the blend and in between P3HT is crystallized (right) into fibre-like structures
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Determining absorptive potential variation in electron beam sensitive specimens using a single energy-filtered bright-field TEM image Somnath Bhattacharyya1, Joerg R. Jinschek2 1. Leibniz-Institut für Oberflächenmodifizierung, 04103 Leipzig, Germany 2. FEI Company, Europe Nanoport, Eindhoven, The Netherlands [email protected] Keywords: TEM, projected potential, electron beam sensitive specimen
The growing interest in determining structure of soft matter (such as organic semiconductors, polymers, biological material, etc.) on the nanometer scale gives rise to focus on transmission electron microscopy (TEM) as one of the investigation tools of choice. Since soft matter samples are often very electron beam sensitive, in most instances they do not allow a higher electron dose. Therefore only a single image can be taken before degeneration. Also, soft matter samples are more or less transparent for the TEM’s high-energy electrons because they are mainly composed of low atomic number (Z) elements. This results in low intensity and low signal-to-noise phase contrast TEM images that are difficult to interpret and to quantify. Here, we present a TEM method to measure variation in the absorptive potential, i.e. the combined variation in atomic density and in chemical composition, in a soft matter specimen directly from the intensity profile in a zero-loss energy-filtered bright-field image at zero defocus. The aim of our work is to obtain an easily-applicable technique that provides information related to structure and composition of the electron beam sensitive soft material retrieved just from a single TEM image in absence of having special arrangements in the microscopes (such as low dose, cryo stage, biprism, etc.). The reliability and effectiveness of the proposed method has been tested both using simulation and experiment. While the use of an objective aperture is advantageous by allowing the high-angle electron scattering to be separated and the chemical information to be retrieved, it also limits the resolution (4.8 Å in the present study depending on the size of the chosen objective aperture). On the other hand, our simulation experiments also indicate that by the use of a well-chosen objective aperture the spherical aberration coefficient of the objective lens (CS) as well as the spatial and temporal coherence envelope function of the contrast transfer function (CTF) do have a negligible effect. As a proof of concept we retrieved absorptive potential across a feature (vesicle) next to the nucleus (large dark feature in the centre of the image in Figure 1a) in the tissue sample of zebrafish gill (for the specific biological details please see ref.[1]). The normalized one-dimensional intensity profile across the dotted region of Figure 1b is shown in Fig. 2a and the retrieved absorptive potential profile (scaled to absolute values) is presented in Fig. 2b. The potential drop inside our feature of interest is of about 0.1eV [2]. This is directly related to differences in atomic density and composition in the specimen, which would not be detectable if the specimen would have
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been ‘stained’ with heavy elements (elements with large Z such as Os, U, …) to enhance the contrast in the TEM image. It is known that the retrieved absorptive potential depends strongly on the local diffraction conditions and thus on dynamical scattering [3]. However, as we are focusing on soft materials that do not have strong diffracting characteristics, dynamic scattering effects are negligible. 1. 2. 3. 4.
L. Karlsson, J. Fish Biol. 23 (1983), p.511. S. Bhattacharyya, J. R. Jinschek, Microscopy Research and Technique, submitted. S. Bhattacharyya, C.T. Koch and M. Ruhle, Ultramicroscopy 106 (2006), p. 525. The authors are extremely thankful to Dr. Christoph Koch for fruitful scientific discussion. Heartfelt thanks are expressed to Sandy Hancock of Virginia-Maryland Regional College of Veterinary Medicine (based at Virginia Tech) for providing the zebrafish gill test sample.
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Figure 1. In-focus (Δf=0) zero-loss energy-filtered bright-field TEM images, (a) overview of the tissue sample of zebrafish gill showing a chloride cell, (b) showing feature of interest (vesicle, indicated by an arrow in (a)) at higher magnification.
Figure 2. (a) Normalized one dimensional intensity profile of the area indicated by dotted box in 1(b), (b) retrieved absorptive potential in absolute scale.
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Elemental distribution of soft materials with newly designed 120kV TEM/STEM C. Hamamoto1, N. Endo1, H. Nishioka1, T. Ishikawa1, Y. Ohkura1, and T. Oikawa2 1. JEOL. Ltd., Electron Optics Division, 1-2 Musashino 3-chome, Akisima, 196-8558, Tokyo, Japan 2. JEOL (EUROPE) SAS, Espace Claude Monet, 1, allee de Giverny 78290 Croissy-sur-Seine, France [email protected] Keywords: STEM、EDS、Soft materials
TEM is commonly used for morphological observation of soft materials (a polymer, a biological specimen etc.), which consist of light elements. These materials show low contrast relative to inorganic materials in TEM images. The contrast of STEM image can be enhanced through analogue and/or digital signal processing. Lower voltage is advantageous for the image contrast. Some of polymers, biomaterials or living organisms contain an inorganic element into their body to develop their functional characters [1,2]. Elemental analysis of local area in the sample is effective for such materials. Therefore, the high contrast morphological observation and analysis (especially elemental maps) have been requested for those fields to clarify their characters. We report STEM and EDS mapping technique for these soft materials by 120kV TEM/STEM. We chose a coating film, which contained the particles of TiO2, Fe2O3, SiO2 as a polymer specimen. The thin film for this specimen was prepared by an ultra microtome. STEM bright and dark images were taken by a 120kV microscope (JEM-1400: JEOL) equipped with a STEM system and an Energy Dispersive X-ray Spectrometer (EDS) (JED-2300T: JEOL). LaB6 emitter was used for this experiment to improve beam current and diameter for STEM and analysis. A STEM-BF image, a STEM-DF image and EDS maps of the coating film are shown Fig. 1. The acquisition time of EDS mapping was about 30 minutes and the number of pixels for an EDS map was 256 x 256 pixels. A chlorine rich domain is confirmed with an EDS map of chlorine with certain reliability. In a STEM-DF of Fig. 1, the domain also shows the high intensity, reflecting the Z contrast, which is one advantage of HAADF imaging. Thus, soft materials can be observed with high contrast and EDS gives reliable evidence of the specimen. 1. B. M. Novak, Advanced Materials, 5, 6 (1993). p.422. 2. K. Rezwan, Q. Z. Chen, J. J. Blaker and A. R. Boccaccini, Biomaterials, 27, 18 (2006), p.3413.
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Figure 1. STEM-BF, STEM-DF image and EDS maps of the coating film
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Preparation of titanate nanotubes and their polymer composites D. Kralova1, N. Neykova1 and M. Slouf1,2 1. Institute of Macromolecular Chemistry AS CR, Heyrovskeho namesti 2, 162 06 Praha 6, Czech Republic 2. Member of Consortium for Research of Nanostructured and Crosslinked Polymeric Materials (CRNCPM) [email protected] Keywords: TiO2, titanate nanotubes, nanocomposite, polyamide
Titanate nanotubes (Ti-NT) are a novel type of nanoparticles based on TiO2. Untill now, Ti-NT have been studied mostly from the point of view of their unique crystalline structure and morphology [1]. Our goal was to prepare well dispersed, dried Ti-NT in high amounts and use nanopowder as filler for polymer nanocomposites [2]. Ti-NT were synthesized by hydrothermal treatment of TiO2 in concentrated NaOH as reported in literature [3]. The suspension was washed with water to pH = 10 and the nanotubes were isolated by freeze-drying. Morphology and structure of Ti-NT were studied by a number of microscopic methods (FESEM, TEM, EDS, SAED). The dried nanopowder of Ti-NT was used in preparation of polymer composites. Composites of polyamide 6 (PA6; Ultramid B5) with TiO2 micro- and nanoparticles and with Ti-NT were prepared by melt mixing at 260 °C. Specimens for testing of mechanical properties were prepared by injection molding. The morphology of composites was inspected by light microscopy (LM). Samples were cut in 5 μm thin sections which were observed in transmitted light. Particles dispersed in the polymer matrix were investigated by TEM. Samples were cut into 50 nm ultrathin sections which were collected on the microscopic copper grid. The samples were studied in bright field mode at 120 kV. The elastic moduli E(MPa) of composite samples were determined by dynamic mechanical analysis (DMA). The composite of PA6 with 5 % of freeze-dried Ti-NT exhibited a 35 % increase in elastic modulus, which was considerably higher than in the case of an analogous composite with commercial TiO2 nanopowder (14 %) or micropowder (9 %). The observed increase in elastic moduli was in agreement with the improvement of filler dispersion in polymer matrix, which was followed by LM and TEM (cf. Figs. 1, 2). 1. 2. 3. 4.
Y.Q. Wang, G.Q. Hu, X.F. Duan, H.L. Sun, Q.K. Xue, Chem. Phys. Lett. 365 (2002) p. 427–431. D. Kralova, E. Pavlova, M. Slouf, R. Kuzel, Mater. Struct. 15, (2008), no. 1, p. 41-45. T. Kasuga, M. Hiramatsu, A. Hoson, T. Sekino, K. Niihara, Adv. Mater., 11, (1999), No.15, p 1307-1311. The authors are indebted for financial support trough grants MSMT 2B06096, GACR 106/06/0761 and GACR 106/06/0044.
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Figure 1. Micrographs of all studied fillers and polymer composites. Average size and dispersion in polymer matrix decreases from TiO2 micropowder to Ti-NT. microparticles SEM micrographs of (a) Ti-NT, (d) TiO2 nanopowder and (g) TiO2 micropowder. LM and TEM micrographs of PA6 filled with (b,c) Ti-NT, (e,f) TiO2 nanopowder and (h,i) TiO2 micropowder.
Figure 2. Elastic moduli of polyamide (PA6) and its 5% composites with TiO2 micropowder (PA6/mTiO2), TiO2 nanopowder (PA6/nTiO2) and titania nanotubes (PA6/Ti-NT). The moduli were determined from DMA.
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Nanometer size wear debris generated from ultrahigh molecular weight polyethylene in vivo M. Lapcikova1, M. Slouf1, J. Dybal1, E. Zolotarevova2, G. Entlicher2, D. Pokorny3, J. Gallo4 and A. Sosna3 1. Institute of Macromolecular Chemistry AS CR, Heyrovskeho namesti 2, 162 06 Praha 6, Czech Republic 2. Faculty of Science, Charles University, Hlavova 8, 128 40 Praha 2, Czech Republic 3. Orthopaedics Clinic of Motol Hospital, V uvalu 84, 156 06 Praha 5, Czech Republic 4. Orthopaedic Clinic of Faculty Hospital Olomouc, I.P. Pavlova 6, 775 20 Olomouc, Czech Republic [email protected] Keywords: ultrahigh molecular weight polyethylene, nanometer size wear debris, morphology of wear particles
Ultrahigh molecular weight polyethylene (UHMWPE) has been used as a bearing material in total joint replacements (TJR) for more than forty years. The success of UHMWPE in TJRs consists in its good biocompatibility, very good friction properties, satisfactory mechanical performance and high wear resistance. Nevertheless, the friction between the polymer and metallic/ceramics components of TJR produces UHMWPE wear particles, which are released from the joint space and cause complex inflammatory reactions leading to aseptic bone loosening [1]. UHMWPE wear particle sizes range from submicrons to several millimeters. Particles below 10 µm exhibit the highest biological activity. Some in vitro wear particles, produced in joint simulators, were shown to be smaller than 0.2 µm. Recently in vitro particles as small as several tens of nanometers were detected [2]. Our study [3] brings the first proof that nano-sized wear particles are produced also in vivo. UHMWPE wear nanoparticles were detected by high-resolution, field emission gun scanning electron microscopy (FESEM) in periprosthetic tissues of two different patients. The purity of the isolated wear nanoparticles was confirmed by energy dispersive analysis of X-rays (EDS) and infrared spectroscopy (IR). The morphology of wear nanoparticles was determined by image analysis of FESEM micrographs. The average equivalent diameters of wear particles in the first and the second patient were 18.5 and 21.2 nm, respectively. Nanowear debris could be reliably detected only if the isolation protocol included intensive sonication and if high magnification micrographs were employed. 1. 2. 3. 4.
E. Ingham and J. Fischer, Biomaterials 26 (2005) p. 1271. A.L. Galvin, J.L. Tipper, E. Ingham and J. Fischer, Wear 259 (2005), p. 977. M. Lapcikova, M. Slouf, J. Dybal, E. Zolotarevova, G. Entlicher, D. Pokorny, J. Gallo, A. Sosna, Wear (submitted). We kindly acknowledge financial support through grant 2B06096 (Ministry of Education, Youth and Sports of the Czech Republic).
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Figure 1. Illustrative FESEM micrographs of PC membranes with UHMWPE wear particles of patient #1 (left, H1) and patient #2 (right, H2) at very high magnifications.
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Analysis of nano-composites based on carbon nanoparticles imbedded in polymers Kangbo Lu1, Joachim Loos1, Sourty Erwan2, Dong Tang2 1.Laboratory of Materials and Interface Chemistry and Soft Matter Cryo-TEM Research Unit, Eindhoven University of Technology, 5600 MB Eindhoven, the Netherlands. 2. FEI Company, Achtseweg Noord 5, 5600 KA Eindhoven/Acht, the Netherlands. [email protected] Keywords: STEM, electron tomography, quantification, polymer nanocomposites
Nowadays, polymer nano-composite is one of the most interesting research topics of polymer science and materials engineering, especially, when the nano-filler can bring polymer matrix special multifunctional properties. In this case, carbon nano-fillers such as carbon nanotubes (CNTs) and carbon black(CB) are excellent candidates for manufacturing high performance and conductive polymer nano-composites due to their unique properties. Controlling the distribution of carbon nano-fillers in polymer matrix is the key point of improving the material properties. Conventional Transmission Electron Microscopy (CTEM) is the main tool to investigate nanofillers’ distribution. Recently, we have introduced High-Angle Annular Dark Field (HAADF) Scanning Transmission Electron Microscopy (STEM) as a versatile tool for investigation of purely carbon-based functional polymer systems [1]. Due to contrast and sharpness enhancement in HAADF-STEM imaging, morphology details are revealed that are not observable or not as clear in CTEM (Figure 1). Main origin for the contrast achieved in STEM is the density difference between the components of the polymer systems under investigation. As an additional issue, changing the camera length, and hence the minimum scattering angle collected on the HAADF detector, is a way to dosing diffraction contrast in carbon-based crystalline materials. Commonly, polymer materials are electron beam sensitive. In this respect, one advantage of STEM is that the electron dose is low when compared with CTEM. However, in CTEM low dose operation modes are implemented that allow ultimate reduction of the actual dose and thus apply doses that are orders of magnitude lower than for STEM. As consequence, electron beam damage is a critical issue that has to be addressed when applying STEM on polymer systems. Another aspect we like to discuss is the application of tomography for better understanding the local organisation of nano-composites. Commonly, two-dimensional (2D) images are used to provide structure information of the sample. However, 2D images represent a projection of the three-dimensional (3D) volume of the sample, which has a thickness of about 100 nm, which may cause that separated MWCNTs or well distributed CB seem to overlap in the projection. Thus 3D volume information is very helpful in our study to better understand nano-scale organization of polymer composites [2].
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The 3D structure and morphology information can be acquired after reconstruction a series of 2D projections – acquired for both CTEM and HAADF-STEM – which were obtained by tilting the specimen around the tilt axis. The 3D reconstruction of CB/PE composites cross sections are shown in Figure 2. CB is good distributed and forms a conductive network in the polymer matrix. We critically discuss advantages and drawbacks of both CTEM and HAADF-STEM tomography for obtaining quantitative volume data. 1. 2. 3.
Erwan Sourty, Svetlana van Bavel, Kangbo Lu, Ralph Guerra, Georg Bar, Joachim Loos, accepted Ultramicroscopy J. Yu, K. Lu, E.D. Sourty, N. Grossiord, C.E. Koning, J. Loos, Carbon 45, 2897-2903, (2007). The authors would like to thank Erwan Sourty for his help with HAADF-STEM data interpretation. Further, we like to thank Ralph Guerra, Georg Bar and Bob Vastenhout from The Dow Chemical Company, Dow Olefinverbund GmbH, and Dow Benelux B.V., respectively, for their help with the CB/polymer nanocomposite materials. The work forms a part of the Dutch Polymer Institute (DPI) program on quantify polymer nano-composites.
Figure 1. Carbon black embedded in polymer, CTEM (left) vs. HAADF-STEM (right).
Figure 2. Snapshot from 3D reconstruction of the carbon black (gray) imbedded in the polymer matrix.
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New developments in SEM for in situ tensile tests on polymers P. Jornsanoh1, G. Thollet1, C. Gauthier1 and K. Masenelli-Varlot1 1. Université de Lyon, MATEIS UMR 5510, INSA-Lyon, 7 avenue Jean Capelle, 69621 Villeurbanne, France [email protected] Keywords: tensile tests, STEM, SEM, ESEM
Investigations of the microscopic deformation and damage mechanisms give access to very useful information to understand the macroscopic mechanical behaviour of materials. However for non conductive materials, specimens need to be coated with a conductive layer and in some cases the specimen surface has to be chemically prepared to reveal a contrast between the different phases. In situ mechanical testing in a SEM of these materials has been limited to small strain deformation since at larger strain, cracks in the conductive coating induce charging effect and thus hinders the observation of the samples. Moreover in any case and whatever the detection mode - i.e. backscattered electrons (BSE) or secondary electrons (SE) - SEM restricts so far to observations of the sample surface. Conversely to conventional SEM, controlled pressure SEM enables the investigation of non conductive samples without any coating, due to the presence of gaseous molecules in the microscope chamber. This offers the possibility to perform in situ mechanical tests even on non-conductive samples. Moreover, it has previously been shown, during wet-STEM experiments, that the detection of the incident electrons diffused at large angles (HAADF-like mode, further called transmission mode), enables the observation of samples up to a few µm thick with a good contrast [1]. This work presents a new procedure for in situ tensile tests in a SEM, using simultaneously classical detectors and a HAADF-like detector. Both imaging systems were used to explore the changes in the structure of a PVC-based nanocomposite. SE images illustrate the crack initiation and propagation through the specimen (Figure1). Images of crack tips obtained in transmission mode (STEM) show microdeformation and cavitation around the filler particles (Figure2A). The STEM contrast makes obvious the presence of filler particles all along crack paths (Figure 2B). From the images, the local deformation can be completely determined using digital image processing based on images correlation that gives access to the whole deformation field in the sample. 1.
A. Bogner, G. Thollet, D. Basset, P. Jouneau, C. Gauthier, Ultramicroscopy 2005, 104, 290.
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Figure 1. SE image of crack initiation on surface defects
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Figure 2. STEM images of a crack tip (a) and filler particles along the crack path (b)
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A study of the spatial distributions of the carbon blacks in polypropylene composites using TEM-Tomography and quantitative image analysis H. Matsumoto1, H. Sugimori2, T. Tanabe2, Y. Fujita3, H. Sano1 and H. Jinnai2 1. Mitsubishi Chemical Gr. Science and Technology Research Center Inc. 1000 Kamoshida, Aobaku, Yokohama-city, 227-8502, JAPAN 2. Dep. of Macromolecular Science and Engineering, Kyoto Inst. of Technology 3. Japan Polypropylene Co., JAPAN [email protected] Keywords: Transmission electron micro-tomography, Quantitative image analysis, Carbon black / polypropylene electron conductive composites
It is well known that nanometer scale carbon materials such as carbon black (CB) and carbon nanotube were applied for filler of electron conductive polymer composites. However, it seems the knowledge of local structure of those networks is limited, but is imperative for understanding the physical mechanism and for controlling those distributions. It is necessary for us to create innovative methods. In this research, 3D imaging based on TEM of CB aggregates surrounding polymer were carried out. And the algorism, quantitative analysis of interconnecting particles in 3D, was developed by our group. The polymer used in this work was CB/iPP system (Japan Polypropylene, Japan). Concentration of CB in iPP is 5 wt %, and its content volume is above the percolation threshold. The TEM used in this work was a 200kV-TEM ( TECNAI G2 F20, FEI Company, USA) which was equipped field emission gun and post column type imaging filter (GATAN, USA). Alignment and reconstruction of tilt images for tomogram were carried out using gold marker tracking method and filter back projection method, respectively. The used software of computed tomography was IMOD imaging software[1]. Figure 1 shows CB aggregations surrounding iPP matrix. From this result, it was impossible to explain an absolute quantity of interconnecting CB particles, because of the projection principle of TEM. Therefore, TEM-T experiment was carried out, and then, quantitative image analysis developed in our group was applied. Figure 2 shows results of quantitative image analysis at the same area of Fig.1. In this image analysis, interconnecting aggregations were classified each other using the packing particles algorithm developed by H. Jinnai [2] after segmentation of reconstruction images, and all interconnecting pathway of CB particles were calculated. From these processes, CB aggregations in Fig.1 could be separated 5. In Fig.2(a), the biggest one was visualized as orange solid circles. In view of electron charge transfer pathway, if it is capable to assume that the pathway of charge transfer is minimum length between one side and the other side of a CB aggregation, the minimum interconnecting pathway from the network was calculated by the algorism of warshall-floyd method [3]. The pathway of each S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 773–774, DOI: 10.1007/978-3-540-85226-1_387, © Springer-Verlag Berlin Heidelberg 2008
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aggregation were visualized as yarrow and red lines in Figure 2 (b), respectively. The pathway of orange one in Figure 2(a) was measured ca. 1.31 μm (indicating a black arrow). 1. 2. 3.
J.R. Kremer, D.N. Mastronarde and J.R. McIntosh, J. Struct. Biol. 116(1996) p.71; D.N. Mastronarde, J. Struct. Biol. 120(1997) p.343; Also://http://bio3d.colorado.edu/imod/. H.Jinnai et al., Macromolecules, 40(2007) p.6758. C. Thomas, L. Charles and R. Ronaid, Introduction of algorithms, first edition. MIT press and McGrawHil, US, (1990).
500 nm Figure 1. Bright field TEM image of aggregated CB surrounding iPP matrix.
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Figure 2. the results of quantitative image analysis of CB aggregates in Fig.1. (a) A segmented digital slice image. Each color-label indicates interconnecting CB aggregates, and (b) a volume rendering image.
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A study of the chain-folded lamellae structure and its array in the isotactic polypropylene spherulites by HAADF-STEM and HV-TEM Tomography techniques H. Matsumoto1, M. Song2, H. Sano1, M. Shimojo2,3 and K. Furuya2 1. Mitsubishi Chemical Gr. Science and Technology Research Center Inc. 1000 Kamoshida-cho, Aoba-ku, Yokohama, 2278502, Japan 2. High-Voltage Electron Microscopy Station, Dep. of Materials Infrastructures, National Inst. for Materials Science, sakura 3-13, Tsukuba, Ibaragi, 3050003 Japan 3. Saitama Inst. of Technology, Fusaiji 1690, Fukaya, Saitama, 3690293, Japan [email protected] Keywords: HAADF-STME and HV-TEM tomography / isotactic polypropylene, 3D array of chain-folded lamellae structures
It seems that the knowledge of the morphologies of chain-folded lamellae structure in the bulk crystalline polymers is limited, however, in industry, is out of necessary to understand increasing mechanical properties. Why had not TEM investigations been advanced? In this answer, two reasons can be thought as follows; 1) Misfit between wide length scales of crystalline polymer, which is classified 3 scales; unit cell (subnanometer), chain-folded lamellae (thickness: 10 nm, lateral dimensions: several um), and finally, spherulites (from μm to mm), and the penetration power of convergent TEM. 2) The lamellae structures in a TEM image are only that of edge-on type in the specimen because of 2D projection principle. In view of wide length scales of lamellae structures, new methods; 3D and high penetration power of HV-TEM observation are very appropriate techniques. In this research, HAADF STEM and HV-TEM tomographic investigation of 3D lamellae structure in isotactic polypropylene spherulites (iPP, Japan Polypropylene Co., Japan) were carried out. The stained specimens for HAADF-STEM and HV-TEM tomography were prepared using the ultra-microtorm-sectioning and the focused-ion-beam techniques, respectively. Those tilt series for tomography were alimented and reconstructed using IMOD software developed by Mastronarde et al [1]. Figure 1 shows a 3D surface rendering image of a cross-section of a spherulite reconstructed from HAADF-STEM tilt series. Multi lamellae structures were visible and classified mother (red) and daughter (green) lamellae structures. Almost the all lamellae structural entities in spherulites were built up by daughter lamellae structure, which were explained growth from mother lamellae by Lotz[2]. However the sample thickness for HAADF-STEM was limited and lateral length of lamellae could not be measured.
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Figure 2 shows bright field HV-TEM images of iPP at different tilt angles. The Rodshape specimen in diameter 1.2 μm could be obtained by the HV-TEM, and lamellae structures were visible at each tilt angles as white lines. In this presentation, the shape of lamellae structure and its array in a iPP spherulites based on HV-TEM tomography will be discussed. 1. 2.
1. J.R. Kremer, D.N. Mastronarde and J.R. McIntosh, J. Struct. Biol. 116(1996) p.71; D.N. Mastronarde, J. Struct. Biol. 120(1997) p.343; Also://http://bio3d.colorado.edu/imod/. 2. B. Lotz, J.C. Wittman: J.Polym.Sci., Ed., 24(1986), p. 1541.
Figure 1 A surface rendering image of lamellae structure in iPP. Mother lamellae and daughter lamellae were identified by the pre-observation of nucleated area. The tiltseries for HAADF-STME tomography was obtained from -75 degrees to +75 degrees, and dual axis method was carried out in order to reduce missing zone.
Figure 2 Bright field HV-TEM images of lamellae structure in iPP. The rod shape specimen was prepared parallel to the redial direction of an iPP spherulete. Mother lamellae was nearly parallel to the rod axis.
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Microstructural analysis of ultra-thin nanocomposite layers fabricated by Cu+ ion implantation in inert polymers G. Di Girolamo1, E. Piscopiello1, M. Massaro1, E. Pesce1, C. Esposito1, L. Tapfer1, M. Vittori Antisari2 ENEA, Dept. Adv. Phys. Technol. and New Materials (FIM), 1.Brindisi Research Center, Strada Statale “Appia” km 713, 72100 Brindisi, Italy 2.Casaccia Research Center, Via Anguillarese 301, 00123 Rome, Italy [email protected] Keywords: TEM, ion implantation, polycarbonate
Ion implantation was used to fabricate ultra-thin nanocomposite subsurface layers in inert polymers for applications in mechanics, optics and electronics [1]. Amorphous polycarbonate substrates were implanted at room temperature with low energy Cu+ ions of 60 keV, at 1 μA/cm2 and with doses in a range from 1x1016 to 1x1017 ions/cm2. The nanocomposite surfaces were investigated by transmission electron microscopy (TEM), X-ray diffraction (XRD), optical absorption spectroscopy and electrical conductivity. Cross-section transmission electron microscopy (XTEM) was used to analyze the microstructure and morphology of the Cu-implanted region. TEM experiments show that nanocrystals are formed at ion doses of 1x1016 ions/cm2 (Figure1). The ionimplanted nanocrystals are located at about 50nm-80nm below the polymer surface, in accordance with TRIM calculations (projected range of 75nm and straggling of 20nm). However, at higher ion doses (5x1016 ions/cm2) a continuous thin nanocrystalline copper films is produced. Figure 2 (b) shows the grazing-incidence XRD patterns (incidence angle = 1°) recorded prior and after Cu-implantation in polycarbonate as well as the corresponding diffraction difference curve. The observed diffraction peaks correspond to copper (cubic phase) in accordance with the ICDD (card no. 851326; JCPDS-ICDD 2000). The XTEM image (Figure 2a) shows a continuous polycrystalline copper films below the polycarbonate surface; the lattice fringes are well observed in the Cu film. Optical absorption spectra show a surface plasmon resonance at 2eV suggesting the formation of nanocrystalline Cu films. This characteristic SPR peak is well pronounced for doses of 5x1016 ions/cm2, while at higher doses the SPR peak is smeared out. This finding is in agreement with the XRD and TEM results that indicate a damaged and structurally disordered film for doses of 1x1017 ions/cm2. In addition, electrical conductivity measurements clearly show a reduced electrical resistance for the samples implanted with a doses of 5x1016 ions/cm2, in accordance with the formation of a continuous metallic film (Figure 2). However, no electrical conductivity could be measured for doses of 1x1016 ions/cm2, since only isolated nanocrystals (no continuous films) are formed. Also higher doses (1x1017 ions/cm2) are detrimental for the electrical properties due to the induced ion radiation damage. 1.
D. Fink, Transport Processes in Ion-Irradiated Polymers, Springer-Verlag (2004)
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 777–778, DOI: 10.1007/978-3-540-85226-1_389, © Springer-Verlag Berlin Heidelberg 2008
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Figure 2. – (a) XTEM image of Cu-implanted polycarbonate showing the formation of a continuous nanocrystalline copper film below the polycarbonate surface. (b) Grazingincidence X-ray diffraction patterns of the polycarbonate prior and after the Cuimplantation; the diffraction difference curve exhibits the characteristic Bragg peaks of the cubic copper.
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In-situ experiments on soft materials in the environmental SEM – Reliable results or merely damage? P. Poelt1, H. Reingruber2, A. Zankel1 and C. Elis1 1. Institute for Electron Microscopy, Graz University of Technology, Steyrerg. 17, A-8010 Graz, Austria 2. Institute of Experimental Physics, Graz University of Technology, Petersg. 16, A-8010 Graz, Austria [email protected] Keywords: environmental SEM, soft materials, irradiation damage
The conventional high vacuum scanning electron microscope is an instrument which is mainly used for the analysis of solid materials. In contrast, the environmental scanning electron microscope (ESEM) is also ideally suited for the operation and control of a great variety of dynamic experiments. Since the type of gas in the specimen chamber, the relative humidity in case of water vapour as gas, the gas pressure and the specimen temperature can be varied over a broad range, an ESEM forms a sort of micro reactor, where the wetting, melting, recrystallization, corrosion … of materials can be investigated. Additionally, no coating of non-conductive materials is necessary to prevent charging. All this seems to make the ESEM an excellent tool for the investigation of soft materials and also their behaviour in a wet environment. But several shortcomings make these experiments much more difficult and the results less reliable than one would predict beforehand. Firstly, soft materials are mainly carbonaceous and therefore give notoriously poor contrast. In the low vacuum the contrast decreases additionally with increasing pressure. Although this could be compensated for by an increase in the probe current, a concurrent increase in the irradiation damage makes this very often impossible [1]. But in many cases a thin coating of the material with e.g. chromium or gold suffices to substantially reduce the damage. Moreover, the presence of water can strongly increase the amount of the irradiation damage due to the formation of highly mobile and reactive free radicals [2, 3]. It can also change the wetting behaviour of the material. Figure 1 shows that both the wetting of a material and its drying-up can be affected by the electron irradiation. But the Figures 2 and 3 prove that despite all these shortcomings the ESEM can be a very valuable tool for in-situ experiments of soft materials. A new and exciting application is automated ultra microtomy in the ESEM and the 3D-representation of the internal structure of materials [4]. Other applications are for example the fracture behaviour of textile fibres in dependence on the relative humidity in the specimen chamber or the imaging of the transport of fluids through porous media [5]. 1. 2.
G.D. Danilatos, Adv. Electronics and El. Phys. 71 (1988), p. 109. C.P. Royall, B.L. Thiel and A.M. Donald, J. Microsc. 204 (2001), p. 185.
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3. 4. 5. 6.
L.M. Jenkins and A.M. Donald, SCANNING 19 (1997), p. 92. W. Denk and H. Horstmann, PLoS. Biol. 2(11) (2004), e329. H. Reingruber, P. Poelt and B. Holst, Proc. 5th World Congress on Industrial Process Tomography, Bergen, Norway (2007), p. 23. We thank Gatan GmbH and Mr. B. Kraus for making the ultra microtome 3VIEWTM available and Mrs. M. Schaffer for creating the 3D-representation.
Figure 1. Wetting and drying-up of a cellulose nitrate membrane (10 nm Au-coating). Left: Before recording the image, mainly the marked area had been irradiated. Contrary to the surrounding, many of the pores in this area are not filled / fully filled with water (image width: 100 µm). Centre and right: Drying-up is delayed in the irradiated areas (centre) compared to other areas (right). In the irradiated areas strong damage is visible (image width: 42 µm).
Figure 2. 3D-representation of EPR (ethylene propylene rubber) modified iPP (isotactic polypropylene) after a tensile test, stopped at 25% yield and stained with RuO4 (160 slices, slice thickness: 100 nm); black: EPR particles; grey: cracks.
Figure 3. Wetting of a polyethersulfon membrane (nominal pore size: 450 nm) with water in dependence on time (width of the images: 45 µm). Some of the big pores remain partially filled / unfilled. The water was provided by condensation at a Peltier cooling stage.
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Structural studies on V-amylose inclusion complexes J.L. Putaux1, M. Cardoso2, M. Morin1, D. Dupeyre1 and K. Mazeau1 1. Centre de Recherches sur les Macromolécules Végétales, ICMG-CNRS, BP 53, F-38041 Grenoble cedex 9, France 2. Brazilian Synchrotron Light Laboratory, P.O. Box 6192, Campinas, ZIP Code 13083-970, Brazil [email protected] Keywords: amylose, single crystals, inclusion complexes, electron diffraction, HREM, modeling
Amylose, the linear homopolymer of α-D-glucose found in native starch, can be crystallized from dilute solutions by addition of a large variety of organic guests (alcohols, lipids, aromas, etc.) [1,2]. The morphology and structure of the resulting V-type crystals depend on the nature of the complexing molecule and parameters such as degree of polymerization (DP), concentration and crystallization temperature [3]. We have used DP 100 amylose biosynthesized in vitro with phosphorylase [4] to prepare single crystals whose morphology and structure were characterized by scanning and transmission electron microscopy, electron diffraction and molecular modeling [5]. Lamellar V-type complexes prepared in the presence of isopropanol and linalool exhibited a rectangular shape (Figures 2a,b). The electron diffraction patterns recorded perpendicularly to the crystal base plane suggested an orthorhombic unit cell containing 6-fold amylose single helices and guest molecules entrapped inside and/or in-between helices [5,6]. Square single crystals of V-amylose complexed with α-naphthol (Figure 2a) yielded exceptional base-plane electron diffraction patterns, up to a resolution of 0.13 nm (Figure 2b). Lattice images at a resolution of 0.39 nm confirmed that amylose was crystallized as 8-fold single helices, in a tetragonal space group (Figure 2c) [7]. The crystal structure was further investigated by molecular modeling to determine the helical conformation and the location of the α-naphthol guest molecules (Figure 2c). Pseudo-spherocrystals made of lamellar subunits were prepared by recrystallizing concentrated (2 wt% ) amylose solutions. Figure 3 shows two examples of peculiar flower-shaped complexes, crystallized in the presence of quinoline (Figure 3a) and α-naphthol (Figure 3b), respectively. 1. 2. 3. 4. 5. 6. 7.
Y. Yamashita, K. Monobe, J. Appl. Polym. Sci. Part A2 9 (1971), 1471. W. Helbert, Doctoral thesis (1994), Université Joseph Fourier Grenoble I, France. A. Buléon, G. Potocki-Véronèse, J.-L. Putaux, Aust. J. Chem. 60 (2007), 706. S. Kitamura S., Yunokawa H., Mitsuie S., Kuge T. Polym. J. 14 (1982), 93. A. Buléon, M.M. Delage, J. Brisson, H. Chanzy, Int. J. Biol. Macromol. 12 (1990), 25. J. Nuessli, J.-L. Putaux, P. Le Bail, A. Buléon, Int. J. Biol. Macromol. 33 (2003), 227. M.B. Cardoso, J.L. Putaux, Y. Nishiyama, W. Helbert, M. Hÿtch, N.P. Silveira, H. Chanzy, Biomacromolecules 8 (2007), 1319.
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Figure 1. TEM micrographs of single crystals of synthetic amylose complexed with isopropanol (a) and linalool (b); c) base-plane electron diffraction pattern of a Visopropanol amylose single crystal.
Figure 2. Single crystals of synthetic amylose complexed with α-naphthol : a) TEM image of lamellar crystals in plan view ; b) base-plane electron diffraction pattern of one crystal; c) translational average of a HREM image of the crystal lattice recorded along the helical axis. Inset : projection of a tentative molecular model, indicating the position of α-naphthol molecules inside and in-between the 8-fold single helices.
Figure 3. SEM images of pseudo-spherocrystals prepared by recrystallizing synthetic amylose in the presence of quinoline (a) and α-naphthol (b).
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Multilamellar nanoparticles from PS-b-PVME copolymers C. Lefebvre1, J.-L. Putaux2, M. Schappacher1, A. Deffieux1 and R. Borsali1,2 1. Laboratoire de Chimie des Polymères Organiques, CNRS-ENSCPB, 16 Av. Pey Berland, F-33600 Pessac, France 2. Centre de Recherches sur les Macromolécules Végétales, ICMG-CNRS, BP 53, F-38041 Grenoble cedex 9, France [email protected] Keywords: block copolymers, micelles, cryo-TEM, cryo-negative staining
Linear poly(styrene-b-vinylmethylether) (PS-b-PVME) has been synthesized using anionic (PS block) and cationic (PVME block) polymerization [1]. The copolymer was first dissolved in THF, a good solvent for both blocks, and water, a selective solvent of the PVME block, was slowly added. Care was taken to operate at a temperature below the low critical solution temperature (LCST) of PVME in water (about 30°C). The copolymer chains self-assembled and formed micelles. Solutions at concentrations from 0.2 to 1 mg/mL were observed by cryo-transmission electron microscopy (cryo-TEM) using a Philips CM200 'Cryo' microscope [1]. PS56-b-PVME126 (PS volume fraction of 0.44) formed two types of micellar assemblies (Figure 1a). The solutions mostly contained cylinders made of an electron-dense PS core (13 nm) and a hardly visible PVME corona. The tubular aspect of the core suggested that residual THF might be entrapped. Images recorded after cryo-negative staining [2] showed wider wormlike micelles (35 nm), indicating that the embedding stain outlined the corona of the cylinders without penetrating in it (Figure 1b). More remarkable was the presence of multilamellar vesicles made of a varying number of uniformly-spaced concentric layers (Figure 2). PS56-b-PVME126 "onions" consisting of up to 13 layers were observed. By extension of the model proposed to explain the contrast of cylindrical micelles in vitreous ice, we assumed that the electron-dense bilayers (6 nm) of the vesicles were formed by the PS blocks and were regularly spaced by two adjacent PVME coronas. Images recorded with a larger defocus revealed the presence of the outer PVME corona (Figure 2b). Work is in progress to determine if the multilamellar vesicles are built by shearing of larger lamellar assemblies [3,4] or through a layer-by-layer mechanism [5,6]. Such PS-b-PVME "solid onions" may constitute interesting multicompartmented nanovectors for encapsulation and controlled release of active molecules. 1. 2. 3. 4. 5. 6.
C. Lefebvre. Doctoral thesis, Bordeaux University (2007). S. De Carlo, C. El-Bez, C. Alvarez-Rúa, J. Borge, J. Dubochet, J. Struct. Biol. 138 (2002), 216. F. Gauffre and D. Roux, Langmuir 15 (1999), 3738. O. Regev and F. Guillemet, Langmuir, 15 (1999), 4357. M. R. Talingting, P. Munk, S. E. Webber, Z. Tuzar, Macromolecules 32 (1999), 1593. H. Shen and A. Eisenberg, Angew. Chem. Int. Ed. 39 (2000), 3310.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 783–784, DOI: 10.1007/978-3-540-85226-1_392, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. Cryo-TEM (a) and cryo-negative staining (b) images of wormlike micelles formed by PS56-b-PVME126 in water.
Figure 2. Cryo-TEM images of multilamellar particles formed by PS56-b-PVME126 in water. The particles are made of 1 and 2 (a), 4 (b) and 6 (c) concentric vesicles.
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TEM/SEM characterisation of hybrid titanoniobiates used as fillers for thermoplastic nanocomposites R. Retoux1, S. Chausson1, L. Le Pluart2, J.M. Rueff1 and P.A. Jaffres3 1. CRISMAT, UMR CNRS 6508, ENSICAEN, and 2. LCMT, UMR CNRS 6507 6 Bd du Maréchal Juin, 14050 Caen cedex France 3. CEMCA, UMR 6521, Faculté des Sciences et Techniques, Université de Bretagne Occidentale, 6 Av. Le Gorgeu, 29238 Brest France [email protected] Keywords: TEM, STEM SEM, nanocomposite, hybrid layered materials, titanoniobates
Layered-silicates such as montmorillonite (MMT) have been widely studied as nanofillers to improve the physical properties of polymers (strength, thermal and barrier effects) [1]. Thermoplastic nanocomposites have been characterized combining X ray diffraction and Electron Microscopy techniques. Here we present the results obtained on hybrid nanocomposites where a modified layered titanoniobate has been used to fill two types of polymers, polyethylene (PE) and polyamide 12 (PA12) [2,3]. This mineral oxide, parent of KTiNbO5 [4], presents the advantage of having a well-defined structure at an atomic scale compared to layered silicate clays [5]. It also allows obtaining particles with a high degree of purity, leading to an easier characterisation of the layers by XRD and a greater regularity of the hybrid structures. Electron Microscopy is here used to establish structure-properties relationships of the nanocomposites at nanoscale. First we showed that the intercalation of N-alkyl amines in the interlayer space of the pristine titanoniobate KTiNbO5 (Fig. 1) improve the dispersion and the exfoliation of the oxide in the both polymer matrixes (PE and PA12). Second, the type of matrix strongly influences the exfoliation degree of the nanofiller in the hybrid nanocomposite. Contrary to PA12, in PE, the nanofiller presents tactoids ranging from a few sheets to numbered thicker particles made of several sheets. Figure 2 present respectively (a) the partially exfoliated sheets in an apolar PE and the full exfoliation in a polar PA12 (b). The figure 3 shows the SEM and STEM EDS mappings of these sheets showing that during the melt intercalation of the filler in the matrix no chemical modification of the filler occurs. There is no diffusion of the Ti and Nb atoms in the matrix. One of the main aims of this study is to improve properties like, here for example, thermomechanical properties. This is illustrated in figure 4 for PE and PA12 filled with pristine KTiNbO5 modified by octadecylamine compared to the neat polymers. 1. 2. 3. 4. 5.
A. Okada, O. Kawasumi, A. Usuki, Y. Kojima et al., Mater. Res. Soc. Proc. 171 (1990). S. Chausson, V. Caignaert, R. Retoux, J.M. Rueff et al l, Polymer, 49, 2, (2008), p. 488. S. Chausson, R. Retoux, J.M. Rueff, L. Le Pluart, and P.A. Jaffres, to be submitted. A. Grandin, M.M. Borel and B. Raveau, Journal of Solid State Chemistry 60 (1985), p. 366. A. Beigbeider, S. Bruzaud, P. Médéric, T. Aubry, Y. Grohens, Polymer, 46 (2005), p. 12279.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 785–786, DOI: 10.1007/978-3-540-85226-1_393, © Springer-Verlag Berlin Heidelberg 2008
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(a) (b) Figure 1. SEM images showing the increase of the interlayer space in the amine intercalated oxide (a) compared to the pristine compound KTiNbO5 (b) . 50 nm
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Phase transitions and ordering in liquid crystals – a case study A.K. Schaper Material Sciences Centre, Philipps University, Hans-Meerwein Str., D-35032 Marburg, Germany [email protected] Keywords: liquid crystals, phase transitions, bond-orientational ordering
The particle density fluctuation of condensed matter is caused by thermal motion of the atoms or molecules, and is influenced by the presence of structural defects and lattice disorder [1]. Therefore, the relationship between the density fluctuation and the x-ray scattering intensity extrapolated toward zero angle may serve as quantitative measure of the state of order and the strength of particular defect populations in solids. Here, x-ray density fluctuation measurements were applied to characterizing the phase transitional behavior of a low-molecular weight liquid crystalline (LC) material. Further information on the transition from the crystalline smectic-B phase of this LC material through an intermediate hexatic-B phase to the liquid-like smectic-A and the isotropic phase has been gained from selected area electron diffraction, in conjunction with light microscopy and calorimetric measurements (Figure 1), as well as using wide angle x-ray scattering [2]. A characteristic feature of the diffraction pattern of the hexatic phase is the angular streaking of the reflections which indicates long-range bond-orientational ordering over macroscopic dimensions, but only short-range positional ordering which propagates through domains a few tens of nanometers in size. The modulation of the diffraction in Figure 1 (IIb) originates from one hexatic monodomain and provides, therefore, directly the magnitude of short-range positional disorder in the SBHex phase. The angular intensity distribution is represented by a χ scan as shown in Figure 2 of a 60° segment, fitting of the experimental data set was accomplished according to the theoretical model by Aharony and coworkers [3]. Structural studies as reported here are of particular interest for LCs built-up in a smectic-like fashion showing a negative dielectric anisotropy where the molecular dipole is perpendicular to the molecular axis, and which contain lateral fluorine substituents. Those LC materials are of importance for optimizing the behavior of LC mixtures, e.g., for the development of vertically aligned (VA) displays [4], or of nanostructured materials with 2D ionic conductivity [5]. 1. 2. 3. 4. 5.
W. Wiegand and W. Ruland, Progr. Coll. & Polym. Sci. 66 (1979), p. 355. A.K, Schaper, S.R.,Zhao and A. Kutoglu, Mater. Sci. Forum 539-543 (2007), p. 3485. A. Aharony, R.J. Birgeneau, J.D. Brock and J.D. Litster, Phys. Rev. Lett. 57 (1986), p. 1012 M. Bremer, M. Klasen-Memmer and K. Tarumi, Adv. Mater. 16 (2004), p. 1882. T. Kato, N. Mizoshita and K. Kishimoto, Angew. Chem. Int. Ed. 45 (2006), p. 38.
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788 6. Acknowledgements are due to S.R. Zhao, A. Kutoglu (Marburg) and U. Richter (Halle) for their experimental and theoretical contributions, and W. Ruland (Marburg) for initiating these studies. I’m grateful to Merck KGaA (Darmstadt) for provision of the LC material.
Figure 1. (Ia-e) Characteristic polarized-light micrographs of CBC-55F at 64, 130, 163, 240, and 303 °C, (II) in situ electron diffraction patterns showing the crystalline SB phase (a), the hexatic SB phase (b), the SA phase (c), the nematic phase (d), and the isotropic phase (e); (III) schematics of the LC structures.
Figure 2. Angular intensity scan of a selected reflection arc of the electron diffraction pattern in Figure 1 (IIb) and calculated fitting curve.
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Study of degradation and regeneration of silicon polymers using cathodoluminescence P. Schauer1, P. Horak1, F. Schauer2, I. Kuritka2 and S. Nespurek3 1. Institute of Scientific Instruments of the AS CR, v.v.i., Kralovopolska 147, CZ-61264 Brno, Czech Republic 2. Polymer Centre, Faculty of Technology, Tomas Bata University in Zlin, TGM sq. 272, CZ-76272 Zlin, Czech Republic 3. Institute of Macromolecular Chemistry of the AS CR, v.v.i., Máchova 7, CZ-12000 Prague, Czech Republic [email protected] Keywords: cathodoluminescence, electron beam degradation, poly[methyl(phenyl)silylene], PMPSi, silicon polymers
Silicon polymers are widely studied materials in recent years, because they show very promising properties for the application in electronics. Their properties such as phase of matter, thermochromism, piezochromism, semiconductivity, photoreactivity, light emission and optical nonlinearity, have become very interesting. Our interest has been focused on the group of polysilanes (often also called polysilylenes), especially on poly[methyl(phenyl)silylene] (PMPhSi), having linear backbone of linked silicon atoms. Cathodoluminescence (CL) degradation and regeneration of PMPhSi are presented in this paper. The studied PMPhSi was prepared by the Wurtz coupling polymerization [1]. The low-molecular weight fractions were extracted with boiling diethyl ether. The layers for the CL measurements were prepared from a toluene solution by casting on quartz disk substrates. For electron beam experiments the PMPhSi specimens were covered with Al sputtered film of 50 nm. A modular CL equipment [2] build in our laboratory was used for the study of PMPhSi. A specimen of PMPhSi was positioned at the face of a light guide, and the CL emission was collected from the substrate side of the specimen. At the CL spectra measurement, the emitted light was guided to the entrance slit of the monochromator, and the PMT was positioned at the exit slit of the monochromator. The equipment was controlled by a PC using the IEEE-488 bus. Spectrally resolved CL intensities of PMPhSi are shown in Fig. 1 and Fig. 2. The typical spectrum of PMPhSi is characterised by the dominating narrow UV band (357 nm) and by a weak broad visible band (420-570 nm). The spectra in Fig. 1 were recorded using continual uniform electron beam irradiation. The measurement was started with a non-irradiated specimen, and CL spectrum recording was repeated in intervals as written in the figure labels. It can be seen that the UV emission decreases rapidly, whereas the visible one is almost unchanged during electron beam irradiation. The spectra in Fig. 2 were recorded during regeneration of the degraded specimen kept in vacuum at the room temperature. It can be seen that UV emission is being partially restored during regeneration. The decrease of the UV emission is either a consequence
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Figure 1. Spectrally resolved CL intensities of PMPhSi during degradation by the electron beam of 10 keV, 160 pA.mm-2. (a) Characteristic UV emission, (b) broad visible emission.
Figure 2. Spectrally resolved CL intensities of PMPhSi during regeneration in vacuum at the room temperature. (a) Characteristic UV emission, (b) broad visible emission. of Si-Si bond scission or of Si-Si bond deformation and deviation from the optimally bonded structure [3]. The Si-Si bond scission, crosslinking or weak bond formation are all possible mechanisms, depending on the excitation conditions. A recovery of CL spectra during regeneration indicates a self-healing effect, resulting in reversibility of changes observed. We kindly acknowledge the support of the project [4]. 1. 2. 3. 4.
X.-H. Zhang, R. West, J. Polym. Sci., Polym. Chem. Ed. 22 (1984), p. 15 P. Schauer, R. Autrata, Fine Mech. Opt. 42 (1997), p. 340. R.D. Miller, J. Michl, Chem. Rev. 89 (1989), p. 1359 Work was supported by the Grant No. IAA100100622 of the Grant Agency of the Academy of Science of the Czech Republic.
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Orthogonal self-assembly of surfactants and hydrogelators: towards new nanostructures M.C.A. Stuart1, A.M.A. Brizard2, E.J. Boekema1 and J.H. van Esch 2 1. Electron Microscopy Group, Groningen Biomolecular Sciences and Biotechnology Institute, University of Groningen, Nijenborgh 4, 9747 AG Groningen, The Netherlands 2. Self-Assembling Systems, Delft Chem Tech, University of Delft, Julianalaan 136, 2628 BL Delft, The Netherlands. [email protected] Keywords: cryo-electron microscopy, multicompartment nanostructures, self-assembly.
Self-assembly of small molecular components holds great promises as a bottom-up approach for nano-objects, but functionality of the resulting nanostructures can by far not compete with the sophisticated systems provided by nature. Surfactants, for instance, can lead to a great diversity of aggregates and mesophases (micelles, vesicles, cubic phases…), but with a level of complexity and functionality that still remains limited. Just like in nature, to increase the level of complexity in self-assembling systems, a straightforward approach consists in making use of multiple components that can display orthogonal self-assembly –i.e. the independent formation of two supramolecular structures each with their own characteristic within a single system. More precisely, we have associated surfactants with low-molecular weight hydrogelators: these molecules, based on cyclohexyl-tris-amino acid, can also selfassemble in one direction through the establishment of H-bonds, leading to the formation of a fiber network and consequently macroscopic gels. Work on mixing behavior of surfactants and various gelators have shown the independent formation of a fibrillar network with encapsulated spherical micelles, Figure 1. In order to produce even more complex nanostructures, this approach has been extended to worm-like micelles that can lead to viscoelastic gels, due to their entanglement. Interestingly, the formation of interpenetrating networks, with original and tuneable rheological properties, has been evidenced by cryo-TEM [1]. Screening of various gelators with vesicle-forming surfactants also revealed that most combinations display orthogonal self assembly, Figure 1. Vesicles were indeed successfully incorporated in a highly responsive network of fibers, without any significant disturbance of these two supramolecular structures. By mean of fluorescent spectroscopy, the stability of these encapsulated vesicles with respect to fusion and leakage has also been investigated. This last example has been exploited to successfully develop liposomes with an encapsulated self-assembling hydrogel (“gellosomes”) [1]. The high responsive character of the gelator makes it very interesting as a mimic of cytoskeleton and it is expected that this new type of nanostructure might be of great interest in drug delivery. 1. A. M. A. Brizard, M. C. A. Stuart, K. J. C. van Bommel, A. Friggeri, M. R. de Jong, J. H. van Esch. Angewandte Chemie int. ed. 47, (2008), 2063.
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Figure. 1 Cyclohexyl-tris-amino acid based hydro-gelator molecules have alternating hydrophilic and hydrophobic regions and can assemble into fibrillar networks by Hbonding (a). Cryo-electron microscopy showed orthogonal self-assembly with surfactant micelles (b), worm-like micelles (c) and phospholipid vesicles (d). Formation of “gellosomes”, vesicles with encapsulated gel network (e). Bar 100 nm.
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Structure of particles formed during Se redox process in aqueous polymer solutions E.I. Suvorova1,2, V.V. Klechkovskaya1, M. Cantoni2, and P.A. Buffat2 1. A.V. Shubnikov Institute of Crystallography RAS, Leninsky pr., 59, Moscow 119333, Russia 2. Centre Interdisciplinaire de Microscopie Electronique, Ecole Polytechnique Fédérale de Lausanne, Station 12, CH-1015, Switzerland [email protected] Keywords: Se/polymer particles, redox process, STEM and EDS.
Nanocomposites based on metal or non-metal nanoparticles dispersed within a polymeric matrix may lead to original applications while combining properties from both the inorganic components and the polymers. We report here the results of a structural and morphological TEM study of particles obtained from aqueous solutions by redox process between selenious H2SeO3 and ascorbic acid and stabilized by different amphiphilic polymers which can deeply influence the structure of particles: - N-poly(vinylpyrrolidone) (PVP), Mw= 2.3⋅104 - poly(2-acrylamido-2-methylpropanesulfonic acid) (PAMPS), Mw= 3.0⋅106 - poly[dimethyl(methacryloyloxyethyl)ammonium] methyl sulfate (PDMAEM), Mw= 9.0⋅106 - polyelectrolyte complexes of sodium dodecyl sulfate with a cationic copolymer of Nvinylpyrrolidone and N,N,N,N-triethyl(methacryloyloxyethyl) ammonium iodide (PEC: SDS-CC), Mw= 1.17⋅105 and 3.85⋅106 - oxyethyl cellulose (OEC), Mw= 1.5⋅105 - PVP (PAMPS, PDMAEM or PEC) and bacterial cellulose Acetobacter xylinum (BCAX) The TEM samples were prepared from fine powders of freeze-dried polymers with embedded Se particles or drops of diluted aqueous colloids deposited onto TEM Cu grids coated with an amorphous carbon thin film. TEM/HRTEM and electron diffraction investigation led first to discover that the spheroidal particles formed during the redox process are always Se/polymer composites. Two particles types coexist (Figure 1a): some neither show any diffraction or fluctuation of absorption contrast inside the particles on bright field images (Figure 1b) nor any evidence of crystalline order on diffraction patterns (Figure 1c). The other ones exhibit domains of 5 nm to a few tens nm in size of monoclinic Se. Thus we consider the first ones as homogeneous dispersion of Se atoms or clusters of a few atoms in polymers and the second ones as agglomerates of Se nanocrystallites in a polymer matrix. In some polymers, composite particles are unstable and they spontaneously evolve to monocrystalline Se particles with time. At contrary, in PVP colloids, the Se/polymer composite-particles were stable for at least the 6 months of investigation. Addition of some BC-AX gel-film leads to expel the polymeric material and crystallization of
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monoclinic Se in faulted monocrystalline (fig. 1d, e) or polycrystalline particles on the day-scale depending on concentration of the gel-film. Similarly addition of BC AX gelfilm to the PEC suspension with SDS/CC=0.5 led to Se crystallization while the Se/polymer composite-particles remained stable in PEC with ratio SDS/CC=4. Composite-particles appear deformed under surface tension forces when they are packed in aggregates of a few particles during evaporation of the solvent. They are highly sensitive to electron irradiation, leading to their total disappearance (by Se sublimation or diffusion?) within 30 s of observation under usual TEM observation conditions. Their irradiation resistance increased with addition of BC AX gel-film. The most stable composite-particles were observed in PVP + AX-BC and in PDMAEM. This effect could probably be explained by the partial loss of organic molecules. EDS microanalysis of a few hundreds particles from different colloids confirmed the presence of carbon inside composite particles. The carbon content was evaluated for composite particles assuming that the contribution of the polymer shell covering them and the supporting carbon film didn't change from that for pure Se particles. Carbon was found to be distributed heterogeneously in the composite particles with domains. The values of the Se/C ratio along the diameter of two particles (top and bottom) in the PVP matrix are shown in the Figure 2. We also used the strong chemical affinity of Ag to Se to reveal this latter by adding Ag nanocrystals to the suspension, after the redox reaction, which induced precipitation of Ag2Se crystallites inside the composite-particles. Observing some 300 crystallites about 5 nm in diameter formed in 90 nm particle lead to conclude that each polymer particle is linked to some 50 Se atoms. The question to know if they are atomically distributed or constitute sub-nanometre clusters is still under investigation.
Figure 2. STEM image and EDS analysis of the Se/C ratio across the particles (c)
Figure 1. TEM images (a) of composite (b, c) and crystalline (d, e) particles in PVP+AX-BC matrix.
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Exploring the 3D organisation of high-performance organic solar cells Svetlana van Bavel1,2, Erwan Sourty1,3, Bert de With1 and Joachim Loos1,2 1. Laboratory of Materials and Interface Chemistry and Soft Matter Cryo-TEM Research Unit, Eindhoven University of Technology, The Netherlands 2. Dutch Polymer Institute, Eindhoven, The Netherlands 3. FEI Company, The Netherlands [email protected] Keywords: morphology, photoactive layer, TEM tomography, HAADF-STEM
In high performance organic solar cells, the photoactive layer consists of a blend of electron donor and electron accepter constituents, a so-called bulk heterojunction. The requirements for the morphology of an efficient photoactive layer are nanoscale phase separation, which provides a large interface area for exciton dissociation, and the existence of continuous pathways for transport of free charge carriers to the appropriate electrodes. In this context, research is now focused on a better understanding of the influence of morphology (particularly, the nanoscale organisation of the phase segregated constituents) on the physical properties of the active layer [1-3]. Because the active layer is deposited from solution, several parameters determine the morphology formation, such as the blending ratio, the solvent used and its evaporation rate, post treatments like annealing, etc. In case of solar cells based on poly(3-hexylthiophene) (P3HT) and a methanofullerene derivative (PCBM), controlled annealing treatment, either at elevated temperature or during slow solvent evaporation, is the crucial step to tune the morphology of the photoactive layer towards the design of high-performance solar cell devices. By applying conventional TEM in combination with electron diffraction and TEM tomography techniques, we demonstrate that controlled annealing treatment results in the formation of nanoscale interpenetrating networks with high crystalline order and favorable concentration gradients of both components through the thickness of the photoactive layer (Figure 1). Such a tailored nanoscale morphology accounts for the considerable increase of the power conversion efficiency in corresponding polymer solar cell devices after annealing. In a system of poly[2-methozy-5-(3’,7’-dimethyloctyloxy)-1,4-phenylene vinylene] (MDMO-PPV) and PCBM, up to 80 weight-% of PCBM is needed to get high performance solar cells. PCBM hardly contributes to light absorption and it is still not completely understood why so much of it should be used. By applying HAADF-STEM (high-angle annular dark field scanning transmission electron microscopy) technique, new details are revealed in the morphological organization of MDMO-PPV/PCBM films that are not observable with conventional bright-field TEM. HAADF-STEM is especially beneficial for (nano-)crystalline materials such as PCBM because in this case diffraction contrast can be additionally enhanced. In the films containing 80 wt.-%
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PCBM (that give rise to rather efficient solar cell devices), HAADF-STEM shows clear domain structures as well as segments interlinking these domains, all together building up a percolating nanoscale network (Figure 2). No such network is observed in case of films with 60 wt.-% PCBM corresponding to less efficient devices. 1. 2. 3.
H. Hoppe and N. S. Sariciftci, J. Mat. Chem. 16 (2006), p. 45. X. Yang, J. Loos, S. C. Veenstra, W. J. H. Verhees, M. M. Wienk, J. M. Kroon, M. A. J. Michels, R. A. J. Janssen, Nano Lett. 5 (2005), p. 579. X. Yang and J. Loos, Macromolecules 40 (2007), p. 1353.
Figure 1. Snapshots of a reconstructed 3D volume of P3HT/PCBM photoactive layers: (A) as spin-coated and (B) thermally annealed. The volume dimensions are about 1700 nm x 1700 nm x 100 nm. After thermal annealing, the P3HT nanowires (shown in white) form a pronounced network all over the film.
(A)
(B)
Figure 2. MDMO-PPV/PCBM films with 80 wt.-% PCBM: (A) bright-field TEM image taken under slight defocus conditions and (B) HAADF-STEM image taken in focus at a camera length of 300 mm where, besides a high signal-to-noise ratio, diffraction contrast adds to the mass-thickness contrast making more obvious the presence of a network interlinking the PCBM-rich domains. The scale bar is 200 nm.
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Morphological study on three kinds of two-dimensional spherulites of PBT T. Yoshioka and M. Tsuji Institute for Chemical Research, Kyoto University, Uji, Kyoto-fu 611-0011, Japan [email protected] Keywords: poly(butylene terephthalate) (PBT), spherulite, polymer crystallization
Poly(butylene terephthalate) (PBT) is one of the aromatic polyesters and an excellent engineering thermoplastic. It was reported for PBT that three different types of spherulite (1. usual-positive type, 2. usual-negative type, and 3. un-usual type in which the Maltese cross pattern does not coincide in direction with the crossed polars) are formed under static condition, depending on the crystallization conditions [1,2]. In isothermal melt-crystallization, when crystallization temperature is above 200°C the resulting spherulites are of “usual-positive type”, and below 200°C the resulting ones are of “un-usual type” [1]. In solution-cast crystallization, when crystallization rate is adequately low the resulting spherulites are of “usual-negative type”, and when adequately high the resulting ones are of “un-usual type” [2]. The major object of this study is to clarify their structural origins, each of which determines the optical property of a PBT spherulite, in particular of a two-dimensional one. Three kinds of specimen thin film, each of which contains one of the three types of two-dimensional PBT spherulite, were prepared by essentially following the preparation methods reported in references [1] and [2]. By electron diffraction (ED), it was confirmed that all these kinds of spherulite are made up of only the α-modification crystals. To obtain the information about unit-cell arrangement in every type of spherulite, the selected-area (SA) ED patterns were obtained from the vicinity of the edge of corresponding spherulite, that is to say the patterns were obtained from a specimen area containing the lamellae which grew in the effectively same direction. The arrangement of unit cell in each type of spherulite was determined from SAED patterns, and the projection of a unit cell with one monomer onto the plane perpendicular to the viewing axis (namely, to the normal of specimen surface) is schematically depicted in Figure 1: Figure 1(a) corresponds to the unusual type, and Figures 1 (b) and (c) correspond to the usual-positive type and usual-negative type, respectively. In Figure 1(a), the direction of projection is parallel to the c-axis. The silhouette of one monomer projected along the c-axis onto the a*-b* plane is well represented with an ellipse enclosing the monomer, as depicted in Figure 1(a). Furthermore, the optical anisotropy of one monomer in this projection should qualitatively be well demonstrated by using this ellipse [2]. The optical anisotropy of one lamellar crystal in this projection should be appreciably approximated by this ellipse. Schultz simulated that, if the major axis of optical ellipsoids within a spherulite is tilted by an angle θ from the radial direction of the spherulite, the resulting Maltese cross pattern is oriented at the angle θ to the crossed polars [3]. Therefore, it is speculated that the optical anisotropy expressed
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by the ellipse, the major axis of which leans by an angle of about 60° (namely, by the angle between (010) and (100) planes) to the growth direction of lamella (namely, to the radial direction of the spherulite), determines the optical property of un-usual type spherulite. _ In both Figures 1 (b) and (c), the incident electron beam direction is the [ ,101] _ direction. The silhouette of one monomer in the projection along [ ,101] can be well represented with an ellipse enclosing the monomer, and then it can be anticipated that optical anisotropy is qualitatively expressed well with this ellipse. In the case of usualpositive type (Figure 1(b)), the direction of major axis of the optical ellipsoid is approximately parallel to the growth direction of lamella (namely, to the radial direction of the spherulite). On the other hand, in the case of usual-negative type (Figure 1(c)), the direction of major axis of the optical ellipsoid is approximately perpendicular to the growth direction of lamella (namely, parallel to the tangential direction of the spherulite). Therefore, it is speculated that, in the usual-positive and usual-negative types, the optical anisotropy expressed with the ellipse in Figures 1 (b) and (c) determines the optical property of spherulite. 1. 2. 3. 4.
R.S. Stein and A. Misra, J. Polym. Sci. Polym. Phys. Ed. 18 (1980), 327-342. E.J. Roche, R.S. Stein and E.L. Thomas, J. Polym. Sci. Polym. Phys. Ed. 18 (1980), 11451158. J.M. Schultz, in “Polymer Crystallization –The Development of Crystalline Order in Thermoplastic Polymers–”, Oxford Univ. Press (2001), p.74. T. Yoshioka, T. Fujimura, N. Manabe, Y. Yokota and M. Tsuji, Polymer 48 (2007), 57805787.
_
[2 ,10]*
(a)
_
[1 ,11]*
[010]*
(b)
(c)
Figure 1. Projections of a unit cell with one monomer within one lamellar crystal of the corresponding spherulite, viewed along the normal of specimen surface: (a) the un-usual type, (b) the usual-positive type, and (c) the usual-negative type. In each of (a)-(c), the growth direction of lamella is vertical, as indicated with respective arrows. Closed and open circles indicate oxygen and carbon atoms, respectively. These figures [4] were drawn with “CrystalDesigner Ver.6.0.3”.
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Solution-Grown Crystals of Optically Active Propene–Carbon Monoxide Copolymer T. Yoshioka1, N. Kosaka2, A. Nakayama3, A.K. Schaper4, W. Massa5, T. Hiyama2, K. Nozaki6 and M. Tsuji1 1. Institute for Chemical Research, Kyoto University, Uji, Kyoto-fu 611-0011, Japan. 2. Department of Material Chemistry, Graduate School of Engineering, Kyoto University, Katsura, Nishikyo-ku, Kyoto 615-8510, Japan. 3. Chemistry and Chemical Biology, Gunma University, Kiryu, Gunma-ken 376-8515, Japan. 4. Material Sciences Centre, Philipps University, Hans-Meerwein-Str., 35032 Marburg, Germany. 5. Department of Chemistry, Philipps University, Hans-Meerwein-Str., 35032 Marburg, Germany. 6. Department of Chemistry and Biotechnology, Graduate School of Engineering, The University of Tokyo, Hongo, Bunkyo-ku, Tokyo 113-8656, Japan. [email protected] Keywords: γ-polyketone, alternating copolymer, lattice constants
Asymmetric alternating copolymerization of propene and carbon monoxide (CO) using late transition metal catalyst gives isotactic and optically active α-methyl-γpolyketone [1]. Here we report the lattice constants of crystal for the polyketone. Poly((S)-1-oxo-2-methylpropylene) (hereafter, γ-polyketone) was prepared by asymmetric alternating copolymerization of propene and CO with Pd(II)-(R,S)BINAPHOS catalyst by the procedure of our previous report [2]. The molecular weight was evaluated by GPC with a polystyrene standard to be Mn = 19300, Mw/Mn = 1.43. Firstly, preparation of single crystals by isothermal crystallization from a 0.01wt% solution in 1-octanol was tried, and an optimum preparation condition was determined: at 105°C for 13 h. Figures 1a and 1b show a typical bright-field TEM image of a multilayered lamellar crystal of γ-polyketone with Pt-Pd shadowing and a schematic representation of the crystal, respectively. An equilateral hexagonal-shaped lamellar crystal is observed, although it is not clearly faceted. The lateral dimension of a crystal was in a range of micrometer scale, and the so-called “lamellar thickness” was about 10nm or less, estimated by Pt-Pd shadowing. Therefore, such a solution-grown crystal of γ-polyketone seems to be a folded-chain single crystal. Figure 1c is a selected-area electron diffraction (SAED) pattern from one whole crystal. The existence of C6 axis and two kinds of mirror planes in the pattern indicates that the γ-polyketone forms crystals of a hexagonal crystal system, being different from the perfectly alternating ethylene-CO polyketone that forms orthorhombic crystals of the α-form with 2(1) helical chains [3]. The hk0 net-pattern of the hexagonal crystal is schematically shown in Figure 1d. The 2-dimensional lattice constants can be estimated to be a=b=1.855 (±0.003) nm and γ=120° on the basis of the reflection rings from aluminum which was vapor-deposited onto the specimen. A 3(1) helical structure, which was proposed for a regio- and stereo-irregular propene-CO polyketone under elongation [4], is suggested to each chain stem in the hexagonal crystal, from the above-mentioned results.
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Next, in order to determine the length of c-axis, preparation of thin fibers which can be directly investigated by TEM was tried by electrospinning (e-spinning). The obtained e-spun fibers were annealed at 130°C for 2 days to increase the perfection of crystallographic packing of chain stems in each crystallite of the fibers. The 3dimensional lattice constants were estimated to be a=b=1.861nm and c=0.897nm (γ=120°) from the annealed fibers, while the ones estimated from the as-e-spun fibres were a=b=1.874nm and c=0.895nm (γ=120°). The value (0.897nm) of c-axis obtained from the annealed e-spun fibers is reasoable in this study, because of the similarity of the value of “a” and “b” with the one estimated from the single crystals. Finally, the isothermally crystallized single crystals were collected as a single-crystal mat by filtration, and their crystallographic information was investigated by X-ray diffraction with incident X-ray being parallel to the mat surface. In the resulting diffraction pattern, no specific reflections, which must be assigned to another crystal modification different from the one identified by TEM, were detected. In other words, all the crystals used for determination of the lattice constants can be regarded as having essentially the same crystal structure. We, therefore, conclude that the crystal system for the solution-grown single crystal of γ-polyketone is hexagonal with the lattice constants of a=b=1.855nm, c=0.897nm and γ=120° (α=β=90°). 1. 2. 3. 4.
K. Nozaki and T. Hiyama, J. Organomet. Chem. 576 (1999), 248. K. Nozaki, et al., J. Am. Chem. Soc. 119 (1997), 12779. V. Grayer, et al., Polymer 36 (1995), 1915. / M. Fujita, et al., Macromolecules 34 (2001), 6147. Y. K. Godovsky, et al., Macromol. Chem. Phys. 200 (1999), 2636. (a)
(b)
{110}
{010}
b {100}
{110}
(c)
(d)
{100}
a
{010} b*
a*
Figure 1. (a) Typical multi-layered single crystal of γ-polyketone grown isothermally from a 0.01wt% solution in 1-octanol at 105°C for 13h. (b) Schematic representation of the crystal shown in (a): in (b), the 6 lateral growth faces are indexed on the basis of (d). (c) SAED pattern corresponding to (a). (d) Schematic drawing of the hk0 net-pattern for (c).
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New insights into ultra-high pressure metamorphism from TEM studies F. Langenhorst1 and A. Escudero1 1. Bayerisches Geoinstitut, University of Bayreuth, D-95440 Bayreuth, Germany [email protected] Keywords: UHP metamorphism, high-pressure minerals, transmission electron microscopy
The discovery of coesite (a high-pressure polymorph of silica) and diamond in ultrahigh pressure metamorphic (UHP) rocks from the Alpine and Kokchetav mountain belts has revolutionized our understanding of the geodynamics and processes in deeply subducted crustal slabs [1,2]. Since then, many new natural occurrences of coesite, diamond and other high-pressure phases (e.g. α-PbO2-structured TiO2) were reported. Transmission electron microscopy (TEM) studies did not only play a key role in the identification of these sometimes nanocrystalline high-pressure phases but also provided essential data on their formation mechanisms and conditions as well as on deformation behaviour and mineral reactions in deep regions of subduction zones. In this TEM study, we have examined coesite in a pyrope quartzite from Dora Maira, Alps [3], diamonds and TiO2 phases in UHP gneisses from the Erzgebirge [4-6]. Though natural coesite from Dora Maira is only preserved as inclusions in garnets and is therefore mostly protected from plastic deformation, it was for the first time possible to characterize its slip systems by TEM. The data are essential to understand or predict the fabrics of silica-rich crustal rocks subducted to depths >100 km. Burgers vectors of dislocations in coesite are [100], [001], and [110] (i.e., a, c, and a+b). The (110) plane could be identified as a slip plane. Small prismatic dislocation loops with Burgers vector [010] are also observed and possibly represent water-related defects. The presence of fluids is also obvious from numerous bubbles occurring on {101} Brazil twin boundaries in surrounding quartz. At great depth, the water was probably dissolved in coesite and was then liberated by the back transformation to quartz occurring during exhumation. Metamorphic diamonds from UHP gneisses of the Erzgebirge occur as inclusions in garnet, too and are surrounded by a number of other phases. TEM-EDX analyses show that the mineral assemblage around diamond is composed of intercalated sheet silicates (potassic and sodic micas, chlorite), anatase, quartz, plagioclase, apatite and other rare earth element phosphates (Figure 1). Since most of these phases are hydrous, it was concluded that diamond formed from supercritical C-O-H fluids [4,5], which reacted with surrounding garnet. In order to resolve the question of the redox reaction that led to the precipitation of elemental carbon as diamond, the iron oxidation state of sheet silicates was measured by electron energy loss spectroscopy (EELS). The sheet silicates are anomalously enriched in Fe3+, suggesting an oxidation of iron, compensated by the reduction of the carbon-bearing precursor (possibly CO2).
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The α-PbO2-structured TiO2 high-pressure phase was recently observed as nanosized slab in a twinned rutile from a diamondiferous gneiss of the Erzgebirge, as well [6]. This TEM study suggests a formation of the TiO2 high-pressure phase by martensitic shear deformation. Our reinvestigation of TiO2 phases from diamondiferous gneiss of the Erzgebirge reveals, however, a completely different microstructure. Rutile crystals contain only few dislocations and are devoid of planar defects. The difference in the observed microstructures may be explained by heterogeneity in deviatoric deformation component. Altogether, the TEM observations suggest that subducted crustal rocks can be exhumed from depths up to 150 km. At this depth, supercritical fluids are liberated by decomposition of volatile-bearing minerals in the subducted slab. These fluids influence the transformation kinetics of high-pressure minerals and result into mineral-forming redox reactions. 1. 2. 3. 4. 5. 6.
C. Chopin, Contrib. Mineral. Petrol. 86 (1984), p. 107-118. N.V. Sobolev, V.S. Shatsky, Nature 343 (1990), p. 742-745. F. Langenhorst, J.P. Poirier, Earth Planet. Science Lett. 203 (2002), p. 793-803. B. Stöckert, J. Duyster, C. Trepmann, H.J. Massonne, Geology 29 (2001), p. 391-394. F. Langenhorst, Mitt. Österr. Miner. Ges. 148 (2003), p. 401-412. Hwang S.L., Shen P., Chu H.T., Yui T.F., Science 288 (2000), p. 321-324.
C
Si
K
Na
P
Ti
Garnet
Diamond
Paragonite
Monazite
Figure 1. STEM image and corresponding element maps of a diamond-bearing inclusion in garnet from the Erzgebirge, Germany.
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Characterization of a (021) twin in coesite using LACBED and precession electron diffraction P. Cordier1 and D. Jacob1 and H.-P. Schertl2 1. Université des Sciences et Technologies de Lille Laboratoire de structure et propriétés de l’état solide – UMR CNRS 8008 59655 Villeneuve d’Ascq Cedex, France 2. Institut für Geologie, Mineralogie und Geophysik, Ruhr-Universität Bochum, D44780 Bochum, Germany [email protected] Keywords: precession electron diffraction, LACBED, twin, dislocations
Coesite is a high-pressure polymorph of silica stable in the pressure range 2.5-9 GPa, which corresponds to a minimum depth in Earth of ca. 90 km. Given the ubiquity of silica at the surface of the Earth, coesite represents a good marker of high-pressure processes. Coesite exhibits a monoclinic symmetry with space group C12/c1. Cell parameters are a = 0.71356, b = 1.23692 and c = 0.71736 nm, with β = 120.34°. While monoclinic in symmetry, the coesite lattice has almost hexagonal dimensions with a chex/ahex ratio of 1.73. It is thus possible to describe the coesite structure within a pseudohexagonal cell. In this study, electron diffraction has been used to characterize a (021)-twin in a metamorphic coesite from Parigi, Dora Maira Massif, Western Alps. Due to the quasihexagonal dimensions of coesite, indexation of spot patterns obtained in conventional diffraction is impossible. Two techniques have been used to characterize this defect: large angle convergent beam electron diffraction (LACBED) and precession electron diffraction (PED). In LACBED, the large number of hkl Bragg lines which can be observed enables the determination of absolute orientations (Figure 1). With PED, the absolute indexation of the patterns is made possible through the possibility of measuring spots intensities (Figure 2). In both cases, the orientation relationships between adjacent parts of the twin are characterized unambiguously. The twin is described as a rotation of 89.94° around the [100] axis of the monoclinic C12/c1 coesite. This microscopic description is fully consistent with original descriptions of twinning in synthetic coesite.
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Figure 1. LACBED patterns taken on one side of the twin (variant A) (a), on the twin area (b) and on the other side of the twin (variant B) (c). In (b), the trace of the twin plane is visible and parallel to (021)A and (02-1)B Bragg lines. The mirror symmetry induced by the twin is clearly seen in the enlarged areas of the patterns in (b) and not present in patterns from either the A or B variants.
Figure 2. Experimental patterns taken on each part of the twin for [110] (a and c) and [101] (b and d) orientations.
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Rubens in the Prado National Museum: analytical characterization of ground layers M.I. Báez1, L. Vidal1, M.D. Gayo2, J. Ramírez-Castellanos3, J.L. Baldonedo4 and A. Rodríguez4 1. Dpto. Pintura-Restauración. Universidad Complutense de Madrid (España) 2. Laboratorio Química. Museo Nacional del Prado (España) 4. Dpto. Química Inorgánica. Universidad Complutense de Madrid (España) 5. Centro Microscopía y Citometría. Universidad Complutense de Madrid (España) [email protected] Keywords: Rubens, scanning electron microscopy, artist materials.
The artistic presence of Rubens in Spain is notable and the National Museum of the Prado has an important number of works his. Some of them were made in Spain during their two stays in the Royal court; others come from Ambers (Belgium), where he used to work usually. The peculiar characteristics in the painting method that he used may additionally be useful in dating or authentication of his artworks. Nevertheless, it is important to note that he painted many canvases during his second visit to Spain in a very short time (1628-29). Furthermore, he always used the painting style of his time, such as, reddish ochre grounds from natural colored earths that have a complex chemical composition, (e.g., different aluminosilicates with different ratio of iron oxides). It is proposed to study the grounds that Rubens used in works executed during his time in Spain and at Ambers. The data thus gathered will serve to compare the materials used with one another and with others of the author’s works, also located at the Prado National Museum, regarding which there are reasonable doubts as to whether they were executed in Spain, during his second travel [1]. For the analysis, estratigraphic sections microsamples have been studied, using scanning electron microscopy (SEM), high transmission electron microscopy (HRTEM) and light microscopy (LM), but here we present only the results of SEM examination. For SEM work, samples are included in epoxy resin and prepared in a thin layer on a sample-holder of the same material; the sections must contain all the particles in the microsample unaltered. The obtained results of the study of several microsamples from the canvas entitled “Filopómenes reconocido por unos ancianos en Megara” (1609) (Ambers) (Figure 1), lead to identify the characteristics and nature of the inorganic materials that were used by Rubens. In this occasion, the date and place of the painting are both perfectly dating. Numerous results of the morphologic analysis of the inorganic materials used by Rubens are discussed; moreover, physic-chemical characteristics and chemical distribution (depending on their granulometry, morphology and electronic density) are shown.
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The SEM images of the micro sample sections reveal a number of common general characteristics: considerable compacting in the internal structure of the ground and highly uniform granulometry (Figure 2). Numerous qualitative and semi-quantitative microanalyses by EDS have been performed to locate the pigments and additives used by the artist. The results shows the main and secondary components, such as aluminosilicates containing iron and magnesium, which seem to be an umber earth (Figure 2.a), mixed with different micas (Figure 2.b) and animal-origin calcium carbonate (chalk) (Figure 2.c). Also regular amounts of lead appear; possibly it is the drying used by the artist. 1. This work has been carried forward with funding from the Ministry of Science and Technology under the National Plan for Scientific Research and Technological Development Projects (R&D) (Ref.: HUM2006-01847/ARTE).
Figure 1. Filopómenes reconocido por unos ancianos en Megara (1609). Oil on canvas. Prado National Museum (Madrid).
Figure 2. Ground layer. SEM backscattering sample. a) Aluminosilicate containing iron and magnesium (*). b) Micas (*). c) Chalk (*).
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Development of the FIB-cryo-SEM approach for the in-situ investigations of the elusive nanostructures in wet geomaterials G. Desbois1, J.L. Urai1 1. Structural Geology, Tectonics and Geomechanics, Geological Institute, RWTH Aachen University, Lochnerstrasse 4-20, 52062 Aachen, Germany [email protected] Keywords: FIB-cryo-SEM, geomaterials, halite, clay, elusive structures, grain boundary, porosity
In fluid-filled porous geomaterials, fluid-rock interactions have important effects on their physical and chemical properties. Though the bulk expression of these properties are relatively well known for a number of geomaterials, the relation between nanostructures and macroproperties are poorly understood for a complete understanding of the fluid-rock interactions. Thus, one of the present challenges in experimental geosciences is to directly characterize the structures of the porous media at the nano scale. However, some geomaterials are so fluids-sensitive that investigations on dried samples, required for conventional electron microscopy imaging, are proscribed. For instance, one new alternative for geosciences is to use the cryo-SEM technology which combines the vitrification of the in-situ fluids to stabilize the microstructures and the SEM imaging at high resolution. In addition, the development of ion milling tools, like FIB, directly embedded into the SEM chamber allows the preparation of high quality polished cross-sections suitable for high-resolution imaging. The FIB-cryo-SEM therefore offers a powerful combination for direct and in-situ investigations of the elusive structures in geomaterials at pore scale. We are developing the use of the FIB-cryo-SEM for the study of halite and clays rocks [1,2,3], which are two very important and widespread geomaterials of which the investigations remain difficult due to their high fluid-sensivity. Halite from salt glaciers is much softer than halite in the deep subsurface, and it deforms to very large strains by solution precipitation creep activated by the small grain size and traces of water in the grain boundaries. The role of the small amounts of water in the grain boundaries during this process is not known in any detail. Yet, the use of the FIB-cryo-SEM is suitable to freeze the deformed grain boundary structures which tend to relax fast after removing the active stress, to overcome the problems of dissolutionrecrystallisation artifacts in grain boundary that occurs on dried samples and to give direct evidence of the fluid distribution. In the long term, this will allow us to test the different models for grain boundary structures in solution-precipitation creep, which have been subject of much controversy for the past twenty years. For clays, the morphology of the porosity has a strong effect on many mechanical and transport properties, but its characterization has been mostly indirect until now. On one hand, none of conventionnal approaches is able to directly describe the in-situ porosity at the pore scale, they are limited due to the poor quality of the surfaces which make it difficult to observe and interpret the nanostructures. On the other hand, all of S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 807–808, DOI: 10.1007/978-3-540-85226-1_404, © Springer-Verlag Berlin Heidelberg 2008
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the conventional methods require dried samples in which the natural structure of the pores could be damaged due to the desiccation and dehydration of the clay minerals. We have started to study fine grained Boom clays (Belgium) which are of special interest because it is considered as a potential host formation for the geological disposal of highlevel and long-lived radioactive waste. This study will be the basis for models of transport and chemical reactions in fine grained materials. Our first investigations on wet halite and wet clay materials are very promising and show that it is possible to stabilize the in-situ fluids in grain boundaries or pores by rapid cooling, preserve the natural structures at nano scale, use the FIB milling tool for producing high quality polished cross-sections and for serial-sectioning to reconstruct accurately the grain boundary and the pore space networks in 3D. Thus, we have validated the use of the FIB-cryo-SEM technology for the in-situ investigations of the elusive structures in wet geomaterials without any damages or artifacts. This opens a new field of applications in geosciences. 1. 2. 3.
Desbois G. and Urai J.L. (In submission). In-situ morphology of the meso-porosity in Boom clay (Mol site, Belgium) inferred by the innovative FIB-cryo-SEM method. Geology. Desbois G., Urai J.L., Burkhardt C., Drury M.R., Hayles M. and Humbel B. (2008). Cryogenic vitrification and 3D serial sectioning using high resolution cryo-FIB SEM technology for brine-filled grain boundaries in halite: first results. Geofluids 8 (1), 60–72. Schenk O., Urai J.L. and Piazolo S. (2006). Structure of grain boundaries in wet, synthetic polycrystalline, statically recrystallizing halite – evidence from cryo-SEM observations. Geofluids, 6: 93-104.
Figure 1. I. SEM picture (BSE) of cryo-stabilized brine film in grain boundary close to a triple junction located in a natural polycrystalline salt sample. The surface has been prepared by using the FIB. II. 3D reconstruction of the pore space by FIB serial crosssectioning around a quartz grain in Boom-clay. The thickness of each cross-section is 500 nm. (a) Initial SE pictures and, (b) equivalent segmented pictures.
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TEM applied on the interface characterisation of the replacement reaction chlorapatite by hydroxyapatite U. Golla-Schindler1, A. Engvik2, H. Austrheim3 and A. Putnis1 1. Institute for Mineralogy, University of Muenster, Corrensstr. 24, 48149 Muenster, Germany 2. Geological Survey of Norway, N-7491 Trondheim, Norway 3. Institute for geoscience/PGP, University of Oslo, N-0316 Oslo, Norway [email protected] Keywords: HRTEM, STEM DF, FIB, apatite, replacement reaction
The aim of this work is to understand the mineral replacement mechanism occurring in the transformation from originally Cl-rich apatite and to hydroxyapatite found in south Norway. The hypothesis is that a fluid-mediated mineral replacement mechanism based on interface-coupled dissolution re-precipitation [1] was inducing the phase transformation. For the TEM studies selected apatite grains were chosen, which show nice alteration interfaces (Figure 1a). The TEM samples were prepared with a dual beam microscope (Zeiss CrossBeam 1540EsB). This machine is equipped with a Kleindiek in situ lift out facility and an EsB inlens detector, which enables the detection of the interface in a similar manner than a BSE detector and allows the precise cutting of the TEM Lamella across the interface. The aim was to obtain images of both phases and to yield structural information from the interface between the chlorapatite and the hydroxyapatite. The TEM analysis were performed using two kinds of transmission electron microscopes a JEOL 3010 and a ZEISS LIBRA 200 FE, where conventional diffraction, STEM and high resolution images were obtained. Figure 1b shows a dark field image of the prepared TEM Lamella, where the chlorapatite and hydroxyapatite regions are clearly visible and additionally the sharp interface between both. The electron diffraction pattern Figure 1 c, d, e taken in the chlorapatite, hydroxyapatite region and at the interface on the position 1, 2, 3 in Figure 1b yield identical diffraction pattern except the sharpness of the diffraction spots. The identical orientation and structures for the studied [1-21] zone axis of both phases gives an indication for a topotactical exchange mechanism. This is confirmed by the HRTEM images taken in the chlorapatite and hydroxyapatite region Figure 1f, g that obtained the same lattice spacing. However, there are obvious differences in both images: on the one hand the chlorapatite high resolution image has a lower signal to noise ratio than the hydroxyapatite, and on the other hand the hydroxyapatite has additional contrast related to a small sized porosity. The difference in the quality and sharpness of the diffraction spots and also the difference in the attainable signal to noise ratio of the high resolution images gives the idea that the chlorapatite phase present less long range order and has more destroyed regions, whereas the hydroxyapatite shows a more recovered crystal structure but with
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an additional small sized porosity. All these TEM results fulfill the criteria for an interface-coupled dissolution re-precipitation reaction mechanism. 1.
A. Putnis and C.V. Putnis Solid State Chemistry 180 (2007), p. 1783-1786
Figure 1. TEM analysis of a selected apatite grain. (a) FIB-SEM image with a clear contrast between the chlor- and hydroxyapatite phases (b) STEM dark field image recorded in the LIBRA 200FE. (c-e) diffraction pattern across the interface at position 1, 2, 3 shown in (b) in the chlorapatite phase, on the interface and in the hydroxyapatite phase recorded with the JEOL 3010 (f, g) high resolution images of the chlorapatite and hydroxyapatite phase, respectively.
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Quantitative study of valence states of zirconolites U. Golla-Schindler1, P. Pöml1 1. Institute for Mineralogy, University of Muenster, Corrensstr. 24, 48149 Muenster, Germany [email protected] Keywords: beam damage, valence states, EELS, zirconolite
Minerals that have mixed valence states are widespread and form in many different rock types. On the one hand, the oxidation state can reflect the redox conditions under which the host materials crystallised and is therefore important for answering fundamental questions about Earth’s evolution and structure. On the other hand, in metamorphic and altered rocks, the oxidation state gives important information about the rock forming and alteration processes. The new generation of electron microscopes equipped with an energy filter enables excellent spatial as well as energy resolution, allowing the acquisition of detailed information about the atomic structure, the chemical composition and the local electronic states of the object. This opens new avenues for advanced applications, like establishing the correlation of macroscopic with microscopic and nanoscopic properties in the field of mineralogy. To apply the new facilities and quantitative EELS and ELNES to study these fundamental questions two main problems have to be overcome. These are: artifact-free specimen preparation and the necessity to require spectra free of electron beam damage effects [1,2]. The dose rate has been found to be a decisive factor in enabling the artifact-free study of beam sensitive material. We have shown that with a dose rate of approximately 1.8 x 102 e/nm2s [3] beam damage effects can be avoided for long exposure times and high electron beam doses. For our studies we selected a homogeneous specimen area with t/λ=0.5, resulting in an absolute specimen thickness of approximately 43 nm [4]. The zirconolite mineral system plays an important role in the development of ceramic waste forms (e.g. synroc [5]) for actinides, especially Pu. To analyse the influence of hydrothermal alteration, zirconolites with Ce as an analogue for Pu have been synthesised with varied chemistry. The difference in the ELNES for different valence states shown in Figure 1 can be used for a quantitative study of the valence state of Ce . The ionic radius of Ce3+ and Ce4+ is significantly different therefore it can be expected that they will occupy different crystal sites. The knowledge of the valence state will consequently enable to yield essential information on the site occupancy. The investigations were performed using two different TEM’s. One is a LIBRA 200FE operating at 200 kV, equipped with a field emission gun, a 4 K slow-scan CCD Camera, and a corrected 90° in-column Omega energy filter and the second is the SATEM operating at 200 kV equipped with a monochromator a Cs-Corrector, a corrected 90° in-column Omega energy filter and a 1 K slow-scan CCD camera. S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 811–812, DOI: 10.1007/978-3-540-85226-1_406, © Springer-Verlag Berlin Heidelberg 2008
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1. 2. 3. 4. 5.
L.A.J. Garvie and P.R. Buseck Journal of Physics and Chemistry of solids 60 (1999), p.1943. L.A.J. Garvie et al. American Mineralogist 89 (2004), p.1610. U. Golla-Schindler, R. Hinrichs, P. Pöml., C. Putnis and A. Putnis Quantitative study of valence states of beam sensitive minerals. Microsc. Microanal. 13 Suppl. 2 (2007), p.12661267. R.F. Egerton, Electron Energy-Loss Spectroscopy in the Electron Microscope 2nd ed., Plenum Press, New York, 1996. G.R. Lumpkin, Elements 2 (2006), p.365.
Figure 1. EELS spectra of Monazite and Cerianite. a), b) recorded with the LIBRA 200FE demonstrating the differences in the ELNES for the different valence states Ce3+ and Ce4+. c) recorded with the SATEM showing the improvement due to the enhanced energy resolution and d) a x-ray absorption spectrum for comparison [5].
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Study of Organic Mineralogical Matter by Scanning Probe Microscopy Ye.A. Golubev, O.V. Kovaleva Institute of Geology of Komi SC of RAS, Pervomaiskaya St., 54, 167982, Syktyvkar, Russia [email protected] Keywords: solid bitumens, amber, scanning probe microscopy
Among products of geological processes in the context of occurrence and variety of nanosize structural elements of natural, roentgenoamorphous, organic substances are most interesting objects [1, 2]. Nanosize structures of such substances are named supermolecular structures. In this work, the results of supermolecular structure researches of natural solid bitumens (hydrocarbons) and fossil resins (ambers) are resulted. These substances concern to the most widespread in lithosphere and practically significant of organic mineralogical substances. The supermolecular structures of natural bitumens of the thermal consequent row asphaltites – lower kerites (albertites) – higher kerites (impsonites) – average anthraxolites – higher anthraxolites from the Timan-Pechora petroleum province and Karelian shungite rocks, Russia, were studied in details [2]. Fossil resin samples for our researches transparent grains and grain fragments of the Baltic amber (Kaliningrad region, Russia) with diameter 5-10 cm have used. The used experimental technique were scanning tunneling (STM) and atomic force (AFM) microscopy, following fracture preparation. It should be noted that natural solid bitumens have a mineral multiphase composition. Therefore, the composition of the surfaces under study should be controlled. The analysis of the element distribution on the surfaces under study was performed by an X-ray spectrometer "Link ISIS", combined with SEM JSM6400 (Jeol). Using X-ray spectral analysis, it was shown that mineral impurities were mainly located as scattered inclusions (from one up to several tens of micrometers in size) in a hydrocarbon matrix. So, the nanometer objects found many times on the AFM-images, can be interpreted as bitumen supermolecular structure particles. In this work, we characterized the supermolecular evolution of natural solid bitumens in the carbonization sequence by quantitative parameters. The types of supermolecular structures and sizes of their initial particles have been determined (ex., Fig. 1, a, b). The transfer from fiber structure to globular-fiber structure with the increase of bitumen metamorphism degree from asphaltites up to average anthraxolites has been observed. The sizes of fibers decrease from 250 up to 30 nm from asphaltites up to average anthraxolites. Higher anthraxolites have globular supermolecular structure. It is shown, that the Baltic amber is mainly make-up of incoherent accumulations of densely aggregated globule-like particles of the sizes from 50 to 120 nm (Fig. 1, c). The
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prevailing form of particles is not determined, they represent a row of the forms varying from ellipsoidal globules up to short fibres. Supermolecular particles do not form homogeneous substances, they are associated in various aggregates. 1. 2.
V.F. Pen'kov Genetic mineralogy of carbonaceous substances. Moscow: Nedra Press; (1996), 356 p. Ye.A. Golubev, O.V. Kovaleva, N.P. Yushkin, Fuel, V. 87. (2008). pp. 32–38.
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c)
b)
Figure 1. AFM image of globular structure of shungite’s carbon from Karelia, Russia (a), “brain”-like structure of average anthraxolites fron Lena River, Siberia, Russia (b), globular-fibrous structure of Baltic amber (c).
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Research of Nanoparticle Aggregates from Water Colloidal Solutions of Natural Carbon Substances and Fullerenes by Atomic Force Microscopy Ye.A. Golubev1, N.N. Rozhkova2 1. Institute of Geology of Komi SC of RAS, Pervomaiskaya St., 54, 167982, Syktyvkar, Russia 2. Institute of Geology of Karelian SC of RAS, Pushkinskaya St., 11, 185000, Petrozavodsk, Russia [email protected] Keywords: shungite, carbon nanoparticles, atomic force microscopy
At the present time interest to geological fullerene-like substances grows. The wellknown natural fullerene-like substance is the anthraxolite of shungite rocks (further – shungite) of Karelia, Russia [1]. Interest to research of colloidal solutions of disaggregated shungite substances is determined by an opportunity of the characteristic of aggregation mechanisms of nanoparticles [2]. In addition, films from colloidal solutions represent also independent value as laboratory model of fine-grained geological materials with the peculiar physical and chemical properties testifying to its activation [2]. The films from water colloidal solutions of С60–С70 were investigated in virtue of structural and morphological similarity of structural elements of shungite carbon and fullerene for comparison. In this work the results of studying of morphological features aggregates of carbon nanoparticles, deposited from fine-grained shungite and fullerene water dispersions is carried out by atomic force and electron microscopy. [3]. For formation films the shungite were dispersed by mechanical and ultrasonic means [4]. High-oriented pyrolitic graphite was used as substrates. Drops of suspensions on substrates were drayed. It is shown, that fullerene water dispersions form at drying a thin films from particle aggregates of tens nanometers size. Aggregates form single, double and multijointed chains. Their orientation is chaotic, some microns long. In addition, single particles are observed. Particles can be divided into two types: i) spherical or ellipsoidal globules (Figure 1, a); ii) cup-like particles (Figure 1, b). Globules have height 70 nm, cup-like particles up to 30 nm. The average diameter of particles in fullerene films is 150 nm, distribution similar to normal (Gaussian). Received from shungite water colloids thin film is a set of a multilayered particle deposits, frequently connected and forming "networks" (Figure 2, a). Particles have form of ellipsoidal globules, their average size makes 60 nm. Fullerene films of particle distribution on the sizes is logarithmically normal. The lognormal form of size distribution is typical for aggregates of colloidal particles. The film, received from shungite water dispersions, are generated from units which average size is similar to sizes of carbon globules of shungite rocks (Figure 2, b). Thus,
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the possibility of shungite globule separation is shown by transformation of carbon shungite rocks in water dispersion. 1. 2. 3. 4.
P.R. Buseck, L.P. Galdobina, V.V. Kovalevski, N.N. Rozhkova, J.W. Valley, and A.Z. Zaidenberg, Canadian Mineralogist. V. 35. (1997). pp. 1363–1378. N.N. Rozhkova in “Perspectives of Fullerene Nanotechnology” ed. E. Osawa, (Dordrecht: Kluwer Academic Pub.), (2002), рр. 237-251. Ye.A. Golubev, N.N. Rozhkova, V.N. Filippov, Surface, V. 10. (2007). pp. 47–52. G.V. Andrievsky, V.K. Klochkov, E.L. Karyakina, N.O. Mchedlov-Petrossyan, Chemical Physics Letters. V. 300. (1999). pp. 392–397.
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b)
Figure 1. AFM-images of particles of fullerene aggregates.
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b)
Figure 2. AFM-images of shungite nanostructure (a) and individual shungite globules (b).
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Diffusion in Synthetic Grain Boundaries K. Hartmann1, R. Wirth1, R. Dohmen2, G. Dresen1 and W. Heinrich1 1. GeoForschungsZentrum Potsdam, Section 4.1, Telegrafenberg, 14473 Potsdam, Germany 2. Ruhr-Universität Bochum, Institut für Geologie, Mineralogie und Geophysik, Universitätsstr. 150, 44780 Bochum [email protected] Keywords: Thin film Diffusion, Grain Boundary, Interface
Grain and phase boundaries usually represent only a small volume fraction of a rock. However, their physical and chemical properties strongly influence the macroscopic properties of rocks, such as elasticity, strength, electrical conductivity, and the efficiency of diffusive mass transport. Grain boundary diffusion is normally estimated to be several orders of magnitude higher compared to volume diffusion [1]. Yttrium-Aluminium-Garnet (YAG) bicrystal samples were synthesised for the first time with the wafer direct bonding method [2]. The highly polished and ultra clean crystal surfaces are saturated with pure adsorbed water and are brought into contact at no force. Upon initial contact, hydrogen bonds of the opposing crystal surfaces are expected to form. The adsorbed water readily evaporates at elevated annealing temperatures leaving a synthetic grain boundary behind. Synthetic garnet is used to investigate the grain boundary structure and grain boundary diffusion in a relatively simple system, as a major practical problem with natural materials is the difficulty in controlling their purity as well as stoichiometry. High-Resolution Transmission Electron Microscopy (HREM) and analytical TEM combined with Focussed Ion Beam (FIB) sample preparation was used to investigate the grain boundary structure and its width. Figure 1 shows a straight grain boundary in YAG where the lattice fringes of the two crystals are directly connected, no noncrystalline material is observed. Diffusion experiments are designed in thin-film geometry, such that the grain boundary is perpendicular to the surface covered with the thin-film. Pulsed Laser Deposition (PLD) [3] was used to deposit Nd or Yb doped YAG on the bicrystal. The thin-films were initially amorphous, but during annealing they crystallized using the structure of the bicrystal. Therefore the grain boundary continues within the epitaxially grown thin-film (Figure 2). After diffusion annealing of the bicrystals diffusion profiles were measured with analytical TEM and/or Rutherford Backscattering (RBS). Cherniak [4] measured volume diffusion profiles with Rutherford Backscattering (RBS) of approx. 50 nm after annealing the sample for 2 h at 1300°C. Even though, we choose the same T-tparameters and analytical techniques we could not detect any volume diffusion at all, grain boundary diffusion could not be observed either. After annealing for 17 h at 1300°C, grain boundary diffusion was measured with EDX in TEM, whereas volume diffusion was still undetectable. The diffusion length between the different experimental S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 817–818, DOI: 10.1007/978-3-540-85226-1_409, © Springer-Verlag Berlin Heidelberg 2008
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approaches strongly differs. This may be caused by different defect structures in the substrate, different water activity or a different contact between substrate and source for the diffusing element. More experiments at different temperatures, diffusion times, and thin-film compositions are planned. 1. 2. 3. 4.
Gleiter H., Chalmers B., Progress in Materials Science 16, (1972) p: 77 Heinemann S., Wirth R., Gottschalk M., Dresen G., Physics and Chemistry of Minerals 32 (2005), p: 229 Dohmen, R., Becker, H.-W., Meissner, E., Etzel, T. & Chakraborty, S., European Journal of Mineralogy 14, (2002), p: 1155 Cherniak, D.J., Physics and Chemistry of Minerals 26, (1998), p: 156
Figure 1. Energy-filtered HREM image of the grain boundary in YAG. The inset shows its diffraction pattern, indices are marked with ‘l’ for left ‘r’ for the right crystal site.
Figure 2. Bright field (BF) TEM image of the thin-film diffusion geometry.
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An examination of Van Gogh’s painting grounds using analytical electron microscopy – sem/fib/tem/edx R. Haswell1, U. Zeile2, K. Mensch1 1. Shell Global Solutions International B.V., 1030 BN Amsterdam, The Netherlands 2. Carl-Zeiss NTS GmbH, D-73446 Oberkochen, Germany [email protected] Keywords: TEM, FIB, SEM, EDX, pigments, ground, Van Gogh
In this paper we report the results of an analytical electron microscopy study of the microstructure of the grounds used by Van Gogh. In an initial study we examined samples from three paintings [1] and tentatively concluded that the nature of the barium sulphate used in the grounds was different in each painting. In order to confirm these initial findings we have prepared additional samples from both the original three paintings plus two new ones. The five paintings were all from the French period dating from 1886-1888. The aim of the work was to determine whether the barium sulphate was the same in closely associated works. To this end we have investigated whether the variations in strontium concentration, both between and with-in individual barium sulphate crystals might help answer this question. The inter-particle characterisation of the barium sulphate was made using Scanning Electron Microscopy (SEM) and Energy Dispersive X-ray Spectroscopy (EDX) while Transmission Electron Microscopy (TEM) and EDX was employed for the intra-particle examination. The thin sections for the TEM were prepared using the Focused Ion Beam (FIB) from barium sulphate crystals selected using the SEM/EDX results. Typical SEM backscatter electron (BSE) images from polished cross-section samples from two of the paintings being investigated are shown in Figure 1. In two of the paintings (F377/4 and F546/9) barium sulphate is the only phase present in the ground as illustrated in Figure 1 (a). In the other three paintings the barium sulphate was one of a number phases present in multiple layers of paint. An example is shown in Figure 1 (b). We have found differences in both the overall strontium concentration in the barium sulphate crystals from different paintings as well as variation within a painting itself using SEM/EDX. However the most striking differences were in the intra-particle variation of the FIB sections measured with TEM/EDX using 0.5 μm diameter spots, as shown in Figure 2. These results indicate that - contrary to our initial conclusions - the barium sulphate can actually be divided into two types, namely: type I which has barium sulphate crystals with large intra-particle variation in the strontium concentration (paintings F244/4, F 297 a/2 and F546/9); and type II which has barium sulphate crystals with uniform strontium concentrations although there is a large difference between individual particles (paintings F297/1 and F377/4). These results did not agree with our prior expectations.
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The ability to classify the barium sulphate into different types is being used to add to our knowledge of the methods and materials used by van Gogh, which is helping in the reconstruction of Van Gogh’s oeuvre and attribution. 1. 2.
R. Haswell1 , U. Zeile and K. Mensch, accepted for publication Microchimica Acta, 2008 This work was possible due to the financial support of Shell Netherlands B.V. This work also benefited from fruitful discussions with Ella Hendriks (Van Gogh Museum, Amsterdam) who also helped in the selection of the Van Gogh grounds.
(a)
(b)
atomic Sr concentration normalised to Ba
Figure 1. Figure 1 (a) and (b): SEM backscatter images from paint samples from portrait of Gauguin (F546/9) and basket with pansies (F244/4), respectively. In Figure 1(b) barium sulphate particles are indicated with arrows. 10.0 8.0 6.0 4.0 2.0 0.0 0
1 F244/4
2 F377/4
3 F546/9
4 F297/1
5 F297 a/2 6
Figure 2. The variation in the atomic Sr concentration normalised to Ba measured using TEM/EDX, spot size 0.5 μm’s, from various FIB sections from individual barium sulphate particles from each of the five paintings. Note: the F. numbers used to identify the paintings correspond to the catalogue numbers in Bart de la Faille’s catalogue
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Amorphisation in fresnoite compounds – a combined ELNES and XANES study Th. Höche1, F. Heyroth2, P.A. van Aken3, F. Schrempel4, G.S. Henderson5, and R.I.R. Blyth6 1. Leibniz-Institut für Oberflächenmodifizierung e.V., D-04103 Leipzig, Germany 2. Martin-Luther-Universität Halle-Wittenberg, IZ für Materialwissenschaften, Heinrich-Damerow-Str. 4, D-06120 Halle, Germany 3. Max-Planck-Institut für Metallforschung, Stuttgart Center for Electron Microscopy, Heisenbergstr. 3, D-70569 Stuttgart, Germany 4. Friedrich-Schiller-Universität Jena, Institut für Festkörperphysik, Max-Wien-Platz 1, D-07743 Jena, Germany 5. Department of Geology, University of Toronto, 22 Russell Street, Toronto, M5S 3B1, Canada 6. Canadian Light Source, 101 Perimeter Road, University of Saskatchewan, Saskatoon S7N OX4, Canada [email protected] Keywords: ELNES, XANES, fresnoite
The fresnoite family of minerals (including Ba2TiSi2O8, Ba2TiGe2O8, Sr2TiSi2O8, as well as Ba2VSi2O8, K2V3O8, and Rb2V3O8) has attracted scientific interest not only for its remarkable piezoelectric [1] and optical properties [2] but also due to the occurrence of pentahedrally coordinated Ti4+ and V4+, respectively. L2,3 electron energy-loss near-edge structure (ELNES) and X-ray absorption nearedge structure (XANES) spectra of the latter elements possess particularly wellpronounced peaks due to the narrow natural line width caused by core-hole life-time broadening. While ELNES spectra are typically averaged over a specimen thickness of 50 to 100 nm (but can be excited by a sub-nm probe), two types of XANES spectra, with very different probing depths, are commonly acquired in parallel: total electron yield (TEY) and fluorescence yield (FY) data. TEY XANES spectra probe some 4 nm [3] while FY XANES spectra are estimated to contain information down to a depth of about 50 nm. In the present contribution, we compare XANES spectra of fresnoite compounds recorded at the SGM beamline of the Canadian Light Source (energy resolution below 0.1 eV) with ELNES spectra obtained in a dedicated STEM (VG HB 501) equipped with a cold field-emisison gun and a Gatan Enfina 1000 spectrometer (spectral resolution ~ 0.4 eV). Samples were exposed to ion irradiation of 300 eV, 6 keV, as well as 200 keV. As shown in the cross-sectional TEM micrograph depicted in Fig. 1, the latter acceleration voltage causes superficial amorphisation down to a depth of around 150 nm [4]. The extent of the amorphisation layer decreases with decreasing ion energy and hence the
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various probing depths of ELNES and XANES spectra can be utilised to get deeper insights into coordination changes associated with amorphisation. Moreover, the advantages of enhanced spectral resolution for the investigation of coordination changes are clearly demonstrated. For transition metal-L3 edges, e.g. the Ti-L3 and the V-L3 ELNES, the conclusion is drawn, based on the XANES data, that it is advantageous to study amorphisation processes at a spectral resolution below 100 meV. 1. 2. 3. 4.
S.A. Markgraf, A. Halliyal, A.S. Bhalla, et al. Ferroelectrics 62 (1985) p. 17. Y. Takahashi, K. Kitamura, Y. Benino, et al., Appl. Phys. Lett., 86 (2005) Art.-No. 091110. B.H. Frazer, B. Gilbert, B.R. Sonderegger, G. De Stasio, Surf. Sci., 537 (2003) p. 161. Th. Höche, F. Schrempel, M. Grodzicki, P.A. van Aken, and F. Heyroth, Chem. Mater., 18 (2006), p. 5351.
Ba2TiSi2O8 Glass
Ba2TiSi2O8 Single Crystal 455
460
465
470
Energy Loss [eV]
Figure 1. Ti-L2,3 ELNES spectra of amorphised (by irradiation with 200 keV Ar+) single-crystalline Ba2TiSi2O8 recorded at different depths beneath the surface. For comparison, a spectrum of the identically composed glass is also shown.
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TEM study of Comet Wild 2 pyroxene particles collected during the stardust mission D. Jacob, J. Stodolna and H. Leroux Laboratoire de Structure et Propriétés de l’Etat Solide - UMR CNRS 8008, Université des Sciences et Technologies de Lille – Bât. C6, 59655 Villeneuve d’Ascq, France. [email protected] Keywords: Stardust, TEM, pyroxene
In January 2006, the NASA Stardust spacecraft successively returned to Earth dust from comet 81P/Wild 2, captured in a low-density SiO2 aerogel. Samples of three collected pyroxene-rich particles have been investigated by transmission electron microscopy (TEM). They are coarse-grained Ca-poor pyroxenes with compositions and structures ranging from orthorhombic enstatite to monoclinic pigeonite. The samples originate from terminal particles of two neighbouring tracks (Figure 1). Details about extraction, manipulation and preparation for TEM by ultramicrotomy can by found in [1]. Results were acquired using LaB6 filaments Philips CM30 (300 keV) and FEI Tecnai G2-20 twin (200 kV) microscopes, equipped with Thermo-Noran and EDAX Si-detectors respectively for Energy Dispersive X-ray Spectroscopy (EDX) (see [2] for a full description of the analytical procedure). The general aspect of the ultramicrotomed samples consists of a central part made of crystalline shards, surrounded by a more or less thin and discontinuous rim of dense amorphous SiO2-rich material. Among the three samples, two exhibit very similar and homogeneous compositions and microstructures (C2027,2,69,2,2 and C2027,3,32,2,3). Their composition corresponds to enstatite within the range En94-97Wo2-5Fs2-5. Selected area electron diffraction patterns reveal an orthorhombic Pbca space group. In most of the shards, planar faults parallel to (100) are observed. Lattice fringe images (Figure 2) reveal that they consist in the insertion of one or more clinoenstatite lamellae (fringe spacing ~ 9 Å) in the orthoenstatite matrix (fringe spacing ~ 18 Å). The third sample (C2027,2,69,1,1) is made of clinopyroxene with composition in the range En73-78Wo36Fs18-23. Diffraction patterns indicate a pigeonite monoclinic P121/c1 space group. The dominant microstructure consists in a high density of (100) lamellae (figure 3). Diffraction patterns show that these lamellae are associated with twinned domains. A few chromite exsolutions were detected, in topotactic relationship to the pigeonite host. The sample also contains small olivine grains (Fa21) in inclusion within the pyroxene matrix. In the three samples, dislocations in glide configuration have also been found. In conclusion, the three studied terminal particles are coarse-grained pyroxene which survived to the strong heating associated with the collect (i.e. a full deceleration from 6 km/s along a ~1cm track). They appear relatively undamaged in comparison to the thermally modified grains frequently found in samples extracted from the wall tracks [1, 2]. The microstructure of the studied samples may have been formed by shock deformation, probably prior to the capture into aerogel. Nevertheless the exceptional S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 823–824, DOI: 10.1007/978-3-540-85226-1_412, © Springer-Verlag Berlin Heidelberg 2008
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physical conditions of the collect include a possible intense thermal pulse. A thermal shock as responsible for the observed microstructure cannot be ruled out. 1. 2. 3.
M. E. Zolensly et al, Science 314 (2006), p.1735. H. Leroux et al., Meteoritics & Planet. Sci. (2008), in press. We thank the French Space Agency CNES for their support. Keiko Nakamura Messenger is gratefully acknowledged for the preparation of the ultramicrotomed samples.
Figure 1. Optical photography (credit: NASA/JSC) of the tracks associated with the three studied terminal particles a)
b)
Figure 2. (a) TEM bright-field image of enstatite in C2027,2,69,2,2. The clinoenstatite lamellae parallel to (100) are clearly visible by the 9 Å lattice spacing, whereas the 18 Å lattice spacing corresponds to orthoenstatite. (b) Diffraction pattern. a)
b)
Figure 3. (a) TEM dark-field image showing the (100) twins in pigeonite, sample C2027,2,69,1,1. (b) Diffraction pattern corresponding to the superposition of the P121/c1 [010] and [01 0] zone axes.
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The mechanism of ilmenite leaching during experimental alteration in HCl-solution A. Janßen1, U. Golla-Schindler, A. Putnis 1. WWU Münster, Institut für Mineralogie, Corrensstraße 24, 48149 Münster, Germany [email protected] Keywords: ilmenite, alteration mechanism, dissolution-reprecipitation
Ilmenite (FeTiO3) is an important mineral being the raw material for the production of titanium for the high-tech industries. The production process typically involves acid treatment, which oxidises and removes the Fe, leaving a TiO2-rich phase, generally rutile. In naturally weathered ilmenite, Grey and Reid (1975) first proposed a two-stage alteration mechanism and remains the generally accepted model [1]. In the first stage ilmenite undergoes weathering through oxidation and removal of Fe to form an apparently continuous series of compositions from ilmenite to pseudorutile (ideally Fe2Ti3O9). The Fe is assumed to diffuse out through the unaltered oxygen lattice. Pseudorutile is a transitional phase and undergoes incongruent dissolution to form rutile, hematite and goethite [2]. Understanding the structural and chemical relationships at the nanometre scale between ilmenite to pseudorutile to rutile is essential because it can be help to understand the exact alteration mechanism of ilmenite, and hence optimise the industrial process. A hard rock ilmenite from the Manvers granite pegmatite dike (Canada) with starting composition Fe0.94Mn0.06Ti0.99O3 was used in this study. The mineral was cut into cubes with length of 3 mm. The dissolution experiments were carried out in 0.1 M HCl at 150 °C for 31 days and in 3 M HCl solution at 150 °C for 4 and 5 days. The resulting products were studied by X-ray diffraction, electron microprobe, scanning and transmission electron microscopy. The first results indicate that the alteration proceeds in two distinct stages, each with a sharp interface between the parent phase and the product. The alteration begins at the original ilmenite crystal surface and along cracks through which the fluid can migrate. The first alteration product is pseudorutile – no phases intermediate between ilmenite and pseudorutile were detected. The textural relationship between ilmenite and pseudorutile suggests a coupled dissolution-reprecipitation mechanism rather than a solid-state continuous oxidation and Fe diffusion mechanism. The second stage involves a further dissolution-reprecipitation step to form rutile. Throughout the alteration process the original morphology of the ilmenite is preserved although the product is highly porous. The rutile inherits crystallographic information from the parent ilmenite, resulting in a triply twinned rutile microstructure Figure 1. A fine scale veining in the ilmenite after the experiment was found Figure 2. Characterization of these veins with HRTEM is still in progress. Electron energy loss
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spectroscopy (EELS) will be use to determine the oxidation state of Fe in the starting material and the resulting alteration products. 1. 2.
I.E. Grey and A.F. Reid, The American Mineralogist 60 (1975), p. 898-906. P.A. Schroeder, J.J. Le Golvan and M.F. Roden, American Mineralogist 87 (2002), p. 16161625.
Figure 1. SE – Image of the treated ilmenite surface. Preservation of the crystallographic information: triply twinned rutile after ilmenite.
Figure 2. TEM – image of the vein structure in the ilmenite after the experiment. New crystals grew in the structure .The characterizations of these crystals are still in progress.
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Microstructure and Texture from Experimentally Deformed Hematite Ore K. Kunze1, H. Siemes2, E. Rybacki3, E. Jansen4, H.-G. Brokmeier5 1. Electron Microscopy ETH Zurich (EMEZ), 8093 Zurich, Switzerland 2. Institut f. Mineralogie u. Lagerstättenlehre, RWTH Aachen, 52056 Aachen, Germany 3. Geoforschungszentrum Potsdam, 14473 Potsdam, Germany 4. Mineralogisches Institut, Uni Bonn and FZ Jülich, 52425 Jülich, Germany 5. Institut f. Werkstoffkunde u. Werkstofftechnik, TU Clausthal and FZ Geesthacht, 21502 Geesthacht, Germany [email protected] Keywords: deformation mechanism, dynamic recrystallisation, electron backscatter diffraction, orientation contrast
Relationships between microstructure and texture (crystallographic preferred orientations, CPO) have been reported by several studies on banded hematite ore from Brasil [1, 2]. Pole figure maxima of basal planes are located about normal to foliation, those of prism planes are within the foliation with highest density towards the lineation. This study aims at further understanding of the deformation mechanisms and texture forming processes in experimentally deformed hematite ore. Microstructural observations were performed using reflected light microscopy and SEM orientation contrast imaging, texture measurements were obtained from neutron diffraction [3,4] and from SEM-EBSD orientation mapping [5,6]. Cylindrical samples of fine grained natural hematite ore (diameter 14mm, length 10mm) have been deformed in torsion using a high pressure – high temperature deformation apparatus [7]. Samples were isolated from the pressure medium (argon gas) by a jacket of iron or copper, and separated from this jacket by a thin (0.5mm) Ag-Pd foil in order to minimize the formation of magnetite. Deformation experiments were performed at temperatures between 700°C and 1000°C, a confining pressure of 400MPa, at twist rates corresponding to a maximum shear strain rate of 0.4e-5s-1 and 4.7e-5s-1 , respectively, and to a maximum shear strain of gamma=4.7. After all of the experiments, hematite has dynamically recrystallised and developed a homogeneous polygonal grain fabric with little shape preferred orientation. The average grain size is larger at higher deformation temperature and shows also a gradient across the sample radius, with a remarkable increase near the central axis (Figure 1). The CPO also records a development with increasing shear strain, where an elliptical caxis (0001) maximum forms slightly off the shear plane normal, and where the CPO strength (texture index) increases monotonously with shear strain (Figure 2). The distributions of (11-20) and (10-10) prism poles follow girdles close to the shear plane, with maxima towards the shear direction. It is concluded from the similar microstructures and textures in nature and experiments that the hematite deformed in both cases primarily by dislocation creep accompanied by dynamic recrystallisation.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 827–828, DOI: 10.1007/978-3-540-85226-1_414, © Springer-Verlag Berlin Heidelberg 2008
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C.A. Rosière, H. Siemes, H. Quade, H.-G. Brokmeier and E.M. Jansen, J. Struct. Geol. 23(2001), 1429-1440. J. Bascou, M.I.B. Raposo, A. Vauchez and M. Egydio-Silva, Earth & Planetary Science Letters 198(2002), 77-92. E. Jansen, W. Schäfer and A. Kirfel, J. Struct. Geol. 22(2000), 1559-1564. H.-G. Brokmeier, U. Zink, R. Schnieber and B. Witassek, Materials Science Forum 273275(1998), 277-282 K. Kunze, S.I. Wright, B.L. Adams & D. Dingley, Textures & Microstruct. 20(1993), 41-54. B.L. Adams, S.I. Wright and K. Kunze, Mat. Trans. 24A (1993), 819-831 M.S. Paterson and D. Olgaard, J. Struct. Geol. 22(2000), 1341-1358.
r = 7mm (outside)
(inside) r = 0mm
Figure 1. Orientation mapping of central cut through torsion sample ST27. Color key according to IPF for torsion axis. A gradient of crystal preferred orientations and in grain size evolved from inside (right) to outside (left) of the torsion cylinder.
r = 5mm 4mm 3mm 2mm 1mm 0mm Figure 2. CPO evolution (sample ST04) with distance (r) from torsion axis and therefore with finite shear strain. Pole figure projections onto the shear plane, CPO strength represented by texture index J (random CPO means J=1).
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Identifying pigments in the temple of Seti I in Abydos (Egypt) E. Pavlidou1, H. Marey Mahmoud2, E. Roumeli1, F. Zorba1, K.M. Paraskevopoulos1, M.F. Ali2 1. Physics Department, Aristotle University of Thessaloniki, 54124 Thessaloniki, Greece 2. Conservation Department, Faculty of Archaeology, Cairo University, 12613 Giza, Egypt [email protected] Keywords: Egypt, SEM, FTIR, Pigment
The temple of Seti I in Abydos, a sacred city noted as the most venerated place in Egypt, was built by the 19th dynasty (ca.1294-1279 BCE). The temple is famous for its remarkably unique design; it is in the shape of an “L” and its wall paintings are decorated with the most complete series of Kings and Gods in Egypt, which virtually helped to decode Egyptian history. Our first results concern to samples from these wall paintings which are examined by SEM-EDS and FTIR microscopy in order to identify the used pigments. The dimensions of the samples were about 3x6mm, with blue, green, yellow and red colors on the surfaces. For the FTIR measurements tiny species from the painted surface of the samples were removed and placed on a freshly prepared KBr pellet. The transmittance IR spectra, were obtained with a Perkin-Elmer FTIR microscope, i-series. A database of FTIR spectra from reference materials was used. The above samples along with crosssectioned specimens, were analyzed also by SEM-EDS, using a Jeol 840A Scanning Microscope with an Energy Dispersive Spectrometer attached by Oxford, model ISIS 300. From the EDS analysis of the blue colored surface of the wall-painting specimens (Fig. 1a) are detected Ca (12%), Cu (15%) and Si (33%), while the FTIR spectra that are collected from the same specimens, present characteristic peaks lying mainly between 1280 and 1000cm-1 that are attributed to Si-O-Si stretching vibrations. The comparison of the FTIR spectra (Fig. 2a) with these from our spectral library and the literature [1], leads to the conclusion that the blue color is Egyptian blue (CaCuSi4O10). Additionally, the peak at 1319cm-1 is a strong indication of the presence of calcium oxalate derived from biodegradation process. Studying the green pigment by EDS (Fig. 1b,d) are observed areas with great amounts of Si (43%) and areas with Ca (11%), Cu (13%) and Si (29%). The FTIR spectra from the green specimens (Fig. 2a) are similar with these from blue, presenting mainly a broader peak in the area 850-1250cm-1, indication of a glassy phase. The combination of the above results guide to the conclusion that the used green pigment is Green Frit [2, 3], a material consisted of cuproan wollastonite with large quantity of a glass phase and few bronze residues. Finally the FTIR spectra of yellow and red samples (Fig. 2b) reveal, except of the
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peaks of calcite, the characteristic bands of ochre, which are attributed to aluminosilicate materials such as kaolin. Additionally the presence of Fe (19-30%) in great amounts in EDS analysis (Fig. 1c), affirms the consideration that the used pigments are yellow and red ochre respectively. The pigments are used in thin layers, as it is observed from the optical examination of the specimen and are common in this period of time. 1. 2. 3.
G. A. Mazzocchin, D. Rudello, C. Bragato, F. Agnoli, J. Cul.Her 5 (2004) p. 129 S. Schiegl, K.L. Weiner, A. El Goresy, Naturwissenschaften 76 (1989) p. 393 P. Bianchetti, F. Talarica, M.G. Vigliano, M.F. Ali, J. Cul.Her 1 (2000) p. 179
Figure 1. Cross Section SEM micrographs of blue (a) green (b), red (c) and chemical mapping of green segment (d) 90
80
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50 40 30 20
B lue Green
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R ed R ed O ch re, referen ce
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Figure 2. FTIR blue and green area (a), spectra from red area (b).
1 2 3 4 5
Code AB B2 GR R Y5
Table I. List of analyzed samples Color Materials identification Blue Egyptian blue, calcium oxalate Blue Egyptian blue Green Green frit, gypsum Red Red ochre, calcite, gypsum Yellow Yellow ochre, calcite, gypsum
3500
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Nanostructural study of ground layers of canvas of Rubens at “El Prado” National Museum J. Ramírez-Castellanos1, J.L. Baldonedo2, M.I. Báez3, L. Vidal3, M.D. Gayo4 and M.J. García3 1. Dpto. de Química Inorgánica. Universidad Complutense de Madrid (España). 2. Centro Miscoscopía y Citometría. Universidad Complutense de Madrid (España). 3. Dpto. de Pintura-Restauración. Universidad Complutense de Madrid (España). 4. Laboratorio de Química. Museo Nacional del Prado (España) [email protected] Keywords: Rubens, high resolution transmission electron microscopy, artist materials.
The authors are members of an inter-disciplinary investigation team working in the field of Cultural Heritage Conservation. The aim of this project is the comparative study of grounds of works in the Prado National Museum by Rubens, during his time in Spain and at Ambers (Belgium), due to the peculiar characteristics of the materials used in the coloured grounds -which are of a complex and varied nature- and specifically the peculiarities that they present in Rubens’ canvases, require a detailed nanostructural characterization by means of high-resolution transmission electron microscopy (HRTEM), by using a 300 FEG JEOL electron microscope and an Energy Dispersive Xray Spectroscopy (EDS) microanalysis, in order to determine the nanostructure and chemical composition of the particles forming the pictorial materials [1]. The final properties of crystallized materials depend on different structural and chemical aspects. Furthermore, the presence of defects, crystal size, chemical composition, stechiometry, cationic substitutions and impurities, all lead to chemical and physical property changes. The data thus gathered will serve to compare the materials used with one another and with others of the author’s works, also located at “El Prado” National Museum, regarding which there are reasonable doubts as to whether they were executed in Spain during his second visit (1628-29). This study is made from estratigraphic microsamples taken from works examination object, following a methodology in the preparation that allows maintaining the pictorial layers and as they were applied by the author. This process is complex, because the ultra-thin sections must be stable under the electron beam and, in addition, the sections must contain all the unaltered particles of the microsample. To achieve this, they were included in a suitably fluid, hard after chemical and thermal treatment and chemically neutral Spurr epoxy resin. The ultra-thin sections (50-100 nm thick) were cutted using an ultramicrotome equipped with a diamond knife and a carbon film was evaporated on to their surface [2]. In this contribution, we present the first obtained results corresponding to the work entitled “Filopómenes descubierto por unos ancianos en Megara” by Rubens at Ambers (1609), in collaboration with Frans Snyders (Figure 1). The possible relations
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between the nanostructural features of the used materials and some related aspects, such as origin, elaboration, manufacture, manipulation, etc. will be discussed. Figure 1. Filopómenes reconocido por unos ancianos en Megara (1609). Oil on canvas. Prado National Museum (Madrid). (Sample place is marked as a red circle).
The SAED pattern (Fig. 2a) shows broad diffused scattering and rings at low angles indicate the amorphous nature of the sample (marked as a red circle in Fig. 1). Moreover, the intensity and discrete spots suggest the presence of randomly oriented grains of very small dimensions. The corresponding HRTEM images (Fig. 2b) reveals a complex microstructure, a glassy matrix containing some crystallized domains were found. EDS microanalysis shows the existence of Si, Al, Fe, Mg, Ca, K and Na. In these sense, data seem to confirm that the used grounds by Rubens are mainly composed by Fealuminosilicates, related to feldespast structure. Figure 2. SAED pattern (a) and corresponding HTEM image (b) of the used materials in the coloured grounds.
1. This work has been carried forward with funding from the Ministry of Science and Technology under the National Plan for Scientific Research and Technological Development Projects (R&D) (Ref.: HUM2006-01847/ARTE). 2. M. San Andrés, M.I. Báez, J.L. Baldonedo and C. Barba, Journal of Microscopy 188 (1997), p. 42-50.
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Micro- and nano-diamond particles in carbon spherules found in soil samples Z. Yang1*, D. Schryvers1, W. Rösler2, N. Tarcea3, J. Popp3 1. EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium 2. Institute for Pre- and Early History, University of Mainz, Schillerstrasse 11, D-55116 Mainz, Germany 3. Institute of Physical Chemistry, University of Jena, Helmholtzweg 4, D-07743, Jena, Germany * now at College of Materials Science and Engineering, Hunan University, Changsha, Ch-410082, China [email protected] Keywords: diamond, nanoparticles, microflakes, impact
Carbonaceous spherules of millimeter size diameter and found in the upper soils throughout Europe are investigated by TEM, including SAED, HRTEM and EELS, and Raman spectroscopy. The spherules consist primarily of carbon and have an open celllike internal structure. Most of the carbon appears in an amorphous state, but different morphologies of nano- and microdiamond particles have also been discovered including flake shapes. The latter observation, together with the original findings of some of these spherules in crater-like structures in the landscape and including severely deformed rocks with some spherules being embedded in the fused crust of excavated rocks, points towards unique conditions of origin for these spherules and particles, possibly exogenic [1]. Optical microscopy and SEM reveal mainly cenospheres exhibiting foam-, sponge-, or cell-like internal structures with cell sizes approximately ranging from 10 to 40 micron, as shown in Figure 1. Elemental analyses using EDX show a high portion of C but also considerable amounts of O and no heavy elements. The matrix of the spherules consists of amorphous carbon, with in many cases embedded monocrystalline nanoparticles or defected polycrystalline nanograins, an example of the first shown in Figure 2. Diffraction rings correspond with an fcc-based structure with a lattice parameter of 0.360 nm (adiamond = 0.356 nm). The appearance of the 200 ring, extinct for the perfect diamond structure, can be attributed to the existence of multiple lattice defects in the nanograins or a deviation from the perfect diamond lattice in the nanoparticles. In some specimens, micrometer-sized, flake-shaped diamonds could be identified inside the cell-like structures: an example is shown in Figure 3 together with a set of SAED patterns revealing diamond extinctions in the expected positions. In Figure 4 the characteristic diamond ELNES shape of the C K-edge obtained from such a microflake is shown (the small π* edge originates from amorphous C surface material) together with the plasmon peak at 33 eV, the latter shifting to 24 eV for the nanoparticles. The existence of micrometer sized diamonds in some particles was supported by the observation of the characteristic sharp diamond band at 1332.3 cm-1 in Raman spectroscopy, as shown in Figure 5 [2]. S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 833–834, DOI: 10.1007/978-3-540-85226-1_417, © Springer-Verlag Berlin Heidelberg 2008
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1. 2. 3.
V. Hoffmann, W. Rösler, A. Patzelt, B. Raeymaekers, P. Van Espen, Meteoritics & Planetary Science 40 (2005) A69 Z.Q. Yang, J. Verbeeck, D. Schryvers, N. Tarcea, J. Popp and W. Rösler, (2008) (in press) doi:10.1016/j.diamond.2008.01.104 We kindly acknowledge support of the GOA project on EELS of the University of Antwerp
1.)
2.)
Figure 1. SEM showing foam-like structure in the interior of the spherules. Figure 2. Monocrystalline diamond nanoparticle viewed along <110> cubic zone.
3.) Figure 3. Several diamond microflakes alongside amorphous carbon support together with some SAED patterns with extinctions indicated by crosses in [001] zone.
4.)
5.)
Figure 4. C K-edge ELNES revealing characteristic diamond σ* shape plus plasmon peak at 33 eV in inset. Figure 5. Raman spectrum with diamond peak at 1332.3 cm-1.
835
The use of FIB/TEM for the study of radiation damage in radioactive/non-radioactive mineral assemblages A.-M. Seydoux-Guillaume1, J.-M. Montel1 and R. Wirth2 1. LMTG, UMR 5563 CNRS, UPS, 14 avenue Edouard Belin, 31400 Toulouse, France 2 GFZ, Telegrafenberg, PB 4.1, 14473 Potsdam, Germany [email protected] Keywords: radiation damage, minerals, FIB/TEM
Radiation damage in radioactive minerals has been studied in geosciences for two main reasons. First, U-Th-rich minerals are used for U-Th-Pb datation, and it is essential to understand the effects of radiation damage on lead retentivity. Second, the effect of long term accumulation of radiation damage is a key parameter for assessing the durability of ceramics that could be used for nuclear-waste storage. One strategy typically used is to study naturally radioactive minerals in specific geological contexts by various analytical methods. In contrast to the numerous studies on radiation effects within radioactive minerals, e.g. zircon, monazite, thorite-group…, i.e. “self-damage”, very few have been done on radiation damage effects in Non Radioactive (NR) host minerals. Damage due to irradiation typically appears as concentric structures named "radiohaloes", and are very familiar to petrologists who use them to identify the presence of radioactive minerals in metamorphic or plutonic rocks. Recently, only two papers investigated radiohaloes, in biotite [1] and in chlorite and cordierite [2]. These studies demonstrated that radiohaloes are created by α-particles and correspond only to modifications of optical characteristics of the host mineral. Furthermore, these authors found intensive damage (i.e. amorphous domains visible in the TEM) only in cordierite over a distance of a few tens of nanometers around radioactive inclusions, and assigned them to recoil nuclei. Alternatively the radiohalo may consist of a "large" radioactive (R) /non-radioactive (NR) interface (Figure 1, in Diopside; [3]), between R and NR host mineral, made of completely different minerals. In some case there can also be almost no radiohalo (Figure 1, in Calcite). In this study we present various examples of “radiohaloes” in mineral pairs (thorite/monazite [4], uranothorianite/diopside, uranothorianite /calcite [3]). Most radioactive minerals in rocks are 10-100 µm in size, and the radiohalo thicknesses only 1-30 µm. It is therefore necessary to adapt the techniques to the size of the areas to be investigated. Samples will be characterised by conventional microscopy, SEM, and TEM associated with FIB preparation. This last method is essential because in-situ measurements are needed in order to study the interface between R/NR minerals (Figure 2). 1.
L. Nasdala, M. Wenzel, M. Andrut, R. Wirth, P. Blaum. The nature of radiohaloes in biotite: experimental studies and modeling. American Mineralogist 86, 2001, p. 498-512.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 835–836, DOI: 10.1007/978-3-540-85226-1_418, © Springer-Verlag Berlin Heidelberg 2008
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2. 3. 4.
L. Nasdala, M. Wildner, R. Wirth, N. Groshopf, D.C. Pal, A. Möller. Alpha particle haloes in chlorite and cordierite. Mineralogy and Petrology 86, 2006, p. 1-27. A.-M. Seydoux-Guillaume, J.-M. Montel, R. Wirth, and B. Moine, Radiation damage in diopside and calcite crystals surrounding uranothorianite, in press in Chemical Geology. A.-M. Seydoux-Guillaume, R. Wirth, and J. Ingrin. Contrasting response of ThSiO4 and monazite to natural irradiation. European Journal of Mineralogy 19, 2007, p. 7-14.
Figure 1. Optical microscope image from two examples of radiohaloes in the Tranomaro skarns (Madagascar). Note the presence of cracks around the uranothorianites (UTh) grains within Cpx, and the difference between the radiohaloes in diopside (Cpx) and in calcite (Cc). After [3]
Figure 2. A-C: SEM images of the UTh / Cpx + Cc interface in figure 1A with FIB locations. D-E: TEM images of Cc1-UTh and Cpx-ϕ boundaries. Note the presence of an amorphous phase (~200 nm thick) between Cc1 and UTh. After [3].
837
Non-destructive 3D measurements of sandstone’s internal micro-architecture using high resolution micro-CT E. Van de Casteele1, S. Bugani2, M. Camaiti3, L. Morselli2 and K. Janssens4 1. SkyScan, Belgium 2. Department of Industrial Chemistry and Materials, University of Bologna, Italy 3. CNR – Institute for Conservation and Enhancement of Cultural Heritage, Italy 4. Department of Chemistry, University of Antwerp, Belgium [email protected] Keywords: X-ray micro-CT, sandstone characterization, 3D analysis
Calcareous stones such as Lecce stones have a high porosity which results in a readily uptake of rainwater. Due to the atmospheric pollutants dissolved in the water these stones, used in a lot of historical buildings, are constantly under attack which leads to a decay of the stone [1]. Different kind of organic hydrophobic products such as Paraloid B72 (PB72) and fluorinated rubber (NH) are often applied as protectives with the aim to reduce the corrosion of the material. In order to study the manner in which these treatment products fill the pores a desktop X-ray microtomography system was used. This technique allows the 3D investigation of the internal structure of the stone in a non-destructive way [2,3]. In this research morphological parameters such as the total porosity (as a percentage of the enclosed empty spaces on the volume of interest), pore size distribution, surface-to-volume ratio (which gives an idea of the complexity of the internal structures) and structure model index (SMI) (giving an estimation of the average shape of the pores (0 = ideal plate, 3 = cylinder and 4 = sphere)) were calculated before and after treatment in order to evaluate the changes induced by the polymer application. The 2D reconstructed cross-sections, shown in Figure 1, confirm that Lecce stone has a very complex internal structure. Several different inclusions such as shells with different shapes and sizes (from a few µm up to 1mm, foraminifera in Figure 1) can be clearly distinguished. The 3D rendering of a small portion of the pores network (Figure 2) gives an idea of the complexity and interconnectivity of the internal structure. The pore size distribution (Figure 2) shows that almost 90% of the pores range from 8 to 29µm. The results of the porosity calculation before and after treatment can be found in Table 1. In both cases the variation of the porosity due to the conservation treatments is significant, but very small. The treatments give very high water repellence to the stone, as reported in [4], but they do not drastically change its natural porosity. X-ray micro-CT is a powerful tool for the investigation of the internal structure of sandstone. The reconstructed cross-sections and 3D rendering of the pores network are able to show, qualitatively the shape and quantitatively the dimension of the pores. Moreover, the data processing allows calculating different morphological parameters useful to characterize the stone. Because µCT is a non-destructive technique and it has a high repeatability, the samples can be monitored during the conservation treatments following the changes in porosity of the specimens that may occur.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 837–838, DOI: 10.1007/978-3-540-85226-1_419, © Springer-Verlag Berlin Heidelberg 2008
838
Figure 1. Left: Reconstructed cross-section of a Lecce stone sample scanned with micro-CT at a pixel size of 2.5µm. Right: A zoom of a reconstructed cross-section of a Lecce stone, including the orthogonal views made through the shell in the middle.
Figure 2. Left: 3D rendering of the pore network of a Lecce stone. Right: Pore size distribution calculated with the sphere fitting method [5] Table 1. Porosity calculated before and after the treatment
1. 2. 3. 4. 5.
Product
Before
After
Decrease
PB72
33.1%
29.5%
3.6%
NH
29.4%
26.5%
2.9%
M. Camaiti, S. Bugani, E. Bernardi, L. Morselli and M. Matteini, Applied Geochemistry 22 (2007): p.1248-1254. A. Sasov, Journal of Microscopy, 147(2) (1987): p.169-192. A. Sasov and D. Van Dyck, Journal of Microscopy, 191(2) (1998): p.151-158 S. Bugani, Study of the interactions between nitrogen oxides (NOx) and stone materials treated with conservation products, Master thesis, University of Bologna, Italy, 2004. T. Hildebrand and P. Ruegsegger, Journal of Microscopy, 185 (1997): p.67-75.
Author Index A Abächerli, V. 403 Abe, Y. 59 Abellan, P. 291, 591 Abetz, C. 751 Abetz, V. 751 Aboussaid, K. 235 Abstreiter, G. 295 Abu-Farsakh, H. 83 Acevedo, D. 457 Adams, T.B. 43 Addiego, F. 755 Adikimenakis, A. 55 Adkins, N. 217 Ahrens, B. 551 Aimadeddine, M. 51 Ajroudi, L. 233 Akamatsu, M. 427 Ako, K. 757 Alandes, L. 735 Albrecht, M. 83 Albu, M. 387 Alexe, M. 101, 329 Algarabel, P.A. 607 Algra, R. 159 Ali, M.F. 829 Alloyeau, D. 187 Almeida Filho, A. 401 Almeida, T. 293 Alonso-González, P. 91 Amatucci, G.G. 525 Ambacher, O. 77 Amstatt, B. 85 Andersen, S.J. 395 Andreano, G. 747 Andrews, A. 149 Andrieu, F. 7 Aouine, M. 185 Aouni, A. 75 Arbiol, J. 223, 295, 597, 643
Arenal, R. 117, 167 Arkharova, N. 119 Arnal, V. 51 Arnberg, L. 435 Arnold, B. 719 Arnoldi, F. 411 Arroyo Rojas Dasilva, Y. 637 Arruebo, M. 597 Aschenbrenner, T. 81 Ash, P. 249 Åsholt, P. 435 Attané, J.P. 613 Austrheim, H. 809 Auzely-Velty, R. 745 Ávila, D. 325 Ávila-Brande, D. 173 Aydemir, U. 531 Ayoub, J.P. 677
B Backen, E. 371 Báez, M.I. 805, 831 Bai, X.D. 115 Bailly, A. 151 Baitinger, M. 531 Bakkers, E.P.A.M. 159 Baldonedo, J.L. 805, 831 Ballif, C. 335 Bals, S. 141, 273, 739 Baluc, N. 503 Bamba, G. 415 Bando, Y. 115 Banerjee, S. 495, 641 Banhart, F. 121, 155 Banhart, J. 279 Baram, M. 521 Baratto, C. 127 Barbot, J.-F. 663 Bargar, J. 315
840
Barlas, B. 435 Barna, P.B. 389 Baro, M.D. 311 Baron, T. 125, 151 Barrett, N. 151 Barthlott, W. 743 Bartova, B. 383, 391, 419, 531 Batov, D.V. 179 Battezzati, L. 437 Bauer, M. 15 Bayle-Guillemaud, P. 189, 611, 613 Beck, U. 725 Bécu, L. 757 Beddies, G. 365 Belkadi, A. 639 Bell, A.J. 547 Bellet-Amalric, E. 85 Bellitto, S. 241 Beltrán, A.M. 45, 91 Ben, T. 45, 91 Benaissa, M. 303 Bender, H. 15, 35, 393 Benedetti, A. 327 Benker, N. 285 Bermanec, V. 157 Bernal, S. 183, 239, 271 Bernard, R. 289 Bernier-Latmani, R. 315 Bertagnolli, E. 149 Bertho, S. 759 Bertin, F. 37 Bett, A.W. 669 Bhattacharyya, S. 523, 761 Bidal, G. 33 Bijelić, M. 157 Birajdar, B.I. 329 Biskupek, J. 111, 655 Bittencourt, C. 141 Bjorge, R. 395 Blank, H. 65 Blank, V.D. 175, 179 Błaż, L. 453, 455 Bleck, W. 485 Bleloch, A.L. 305, 331, 365, 587 Blicharski, M. 347
Author Index
Blumtritt, H. 217 Blyth, R.I.R. 821 Bochniak, W. 455 Bocker, C. 523 Boe, A. 509 Boekema, E.J. 791 Boese, M. 325 Boeuf, F. 33 Bogdanoff, P. 279 Bohácek, J. 267 Boldyreva, K. 101 Bonetti, E. 309 Bonnot, A.M. 117, 205 Borgström, M.T. 159 Börjesson, J. 297 Borsali, R. 783 Böttcher, A. 253 Boudin, S. 571 Bougerol, C. 85 Boulanger, L. 397 Boullay, P. 323, 527 Bourgeois, L. 399 Bovi, M. 727 Bowen, J.R. 349 Brabetz, M. 493 Brånemark, R. 741 Brault, J. 303 Bréchet, Y. 427, 483 Briggs, G.A.D. 177 Briot, O. 69 Briston, K.J. 169 Brizard, A.M.A. 791 Brokmeier, H.-G. 827 Brown, A.P. 601 Brown, D.P. 135 Brown, P.D. 113, 293, 731 Browning, N.D. 69 Brun, N. 205, 717 Bruno, P. 167 Brydson, R.M. 215, 601 Buban, J.P. 667 Büchner, B. 307 Buffat, P.A. 315, 681, 793 Bugajski, M. 61 Bugani, S. 837
Author Index
Bullough, T. 143 Bund, A. 685 Burghardt, H. 191 Burnett, T.L. 547 Buso, S.J. 401, 421
C Cabié, M. 193 Cabo, M. 311 Cadete Santos Aires, F.J. 185 Caignaert, V. 323 Caillard, D. 635 Caliste, D. 665 Callini, E. 309 Calmels, L. 631, 715 Calvino, J.J. 183, 199, 213, 235, 239, 271 Camaiti, M. 837 Camassel, J. 57 Campion, R.P. 47 Cantoni, M. 79, 403, 531, 793 Cao, S. 405 Carati, A. 201 Carbó-Argibay, E. 259 Carbone, D. 255 Cardoso, M. 781 Carrillo-Cabrera, W. 407 Casanove, M.J. 291, 591 Casci, J. 249 Castell, M.R. 171 Castell, O. 311 Castro, A. 543 Cauqui, M.A. 213 Cavallotti, P.L. 687 Cedergren, K. 357 Čeh, M. 129, 585 Cellai, L. 747 Cerezo, A 41 Chabli, A. 9 Chaliampalias, D. 701, 703, 705, 707 Chalker, P.R. 143 Chamard, V. 11 Chapman, J.N. 391
841
Chausson, S. 785 Chauvat, M.P. 89 Cheng, C. 723 Cherns, D. 47 Cherns, P.D. 49 Cheynet, M. 33, 51, 195, 683 Chèze, C. 143 Chiba, A. 447 Chisholm, M.F. 45 Chmielowski, R. 347 Chou, Y.H. 215 Chrissafis, K. 701, 703, 705 Chuvilin, A.L. 123, 381 Ciancio, R. 357 Cimalla, V. 77 Cleij, T.J. 759 Clement, L. 9 Clifton, P.H. 41 Cockayne, D.J.H. 177, 519 Coe, S. 731 Coghe, F. 499 Cojocaru, P. 687 Coleman, J.N. 145 Colliex, C. 103, 289, 717 Colombo, C. 295 Comini, E. 127 Comyn, T.P. 547 Cooper, D. 9 Coraux, J. 85 Cordier, P. 803 Cornet, N. 529 Correa-Duarte, M.A. 153 Cosandey, F. 525 Cossange, C. 411 Costa, P.M.F.J. 115 Cotton, N.J. 737 Coulon, P.E. 715 Craven, A.J. 23, 67, 409 Creemer, J.F. 197 Crozier, P.A. 277 Cullis, A.G. 169 Curiotto, S. 437 Czerwinski, A. 61 Czyrska-Filemonowicz, A. 459, 517, 681, 699
842
D Dahl, S. 211 Dahmen, U. 367, 473 Dahoun, A. 755 Daneu, N. 361 Dartsch, H. 81 Daudin, B. 85 Dawson, P. 41 de Dios, S. 567 De Gendt, S. 23 De Keyser, K. 365 De Mierry, P. 71 De Riccardis, M.F. 255 De Teresa, J.M. 607 de With, B. 795 Deak, D.S. 171 Deffieux, A. 783 Delabrouille, F. 411 Delalande, M. 189 Delannay, F. 649 Delaye, V. 7 Deleonibus, S. 7 Delgado, J.J. 199 Delplancke-Ogletree, M.P. 709 Delville, R. 383, 413 Demolon, P. 71 den Hertog, M.I. 125 Denker, C. 93 Denorme, S. 33 Desbois, G. 807 Desinan, S. 271 Desré, P.J. 125 Detavernier, C. 365 Detemple, E. 109 Dey, G.K. 495, 625, 641 D’Haen, J. 759 Dhalluin, F. 125 Di Girolamo, G. 777 Di Martino, J. 755 di Monte, R. 271 Di Paola, E. 201 Díaz-Droguett, D. 203 Dieterle, L. 593 Dietrich, C. 595
Author Index
Dietrich, D. 685, 687 Dietz, W. 729 Dimitrakopulos, G.P. 53, 55, 537, 639, 651 Dimroth, F. 669 Dłużewski, P. 133, 301, 639 Dohmen, R. 817 Dolzhikov, S.V. 471 Dong, C. 319 Donnadieu, P. 11, 415, 417 Dorbandt, I. 279 Dorcet, V. 527 Dörfel, I. 689 Dosch, H. 225 Douthwaite, R.E. 215 Dražić, G. 713 Dresen, G. 817 Dressler, M. 689 Drewello, V. 107 Drube, W. 141 Dubiel, B. 517 Ducati, C. 165 Dudeck, K. 519 Dunin-Borkowski, R. 165 Dupeyre, D. 781 Duppel, V. 653 Duran, A. 523 Durand, D. 757 Dybal, J. 767
E Edmonds, D.V. 429, 431 Edwards, H.K. 731 Eggeler, G. 515 Eibl, O. 351, 353, 355, 371 Eilers, G. 107 Eisenschmidt, C. 551 Elis, C. 779 Emanuelsson, L. 741 Endo, N. 763 Ene, C. 105, 261 Eneman, G. 15 Engel, S. 353, 371 Engelmann, H.J. 13
Author Index
Engqvist, H. 741 Engvik, A. 809 Ensslin, K. 131 Entlicher, G. 767 Epicier, T. 457 Ericson, F. 741 Erni, R. 39, 231, 473 Ernst, T. 7 Ersen, S. 195 Erwan, S. 769 Escribano, S. 207 Escudero, A. 801 Espinoza, R. 419 Esposito, C. 777 Espósito, I.M. 401, 421 Essoumhi, A. 233 Estradé, S. 295, 643 Eswaramoorthy, S.K. 633 Etheridge, J. 229 Eustace, D.A. 177 Ezcurdia, M. 221
F Faglia, G. 127, 747 Falke, M. 331, 365 Falke, U. 331 Fang, Y. 339 Farley, N. 47 Farooq, M.U. 547 Favia, P. 15, 393 Fay, M.W. 113, 731 Faynot, O. 7 Fecht, H. 217 Feiner, L.F. 159 Felten, A. 141 Feltin, E. 79 Fendrych, F. 603 Feng, Y. 305 Fenouiller-Beranger, C. 33 Fernández, A. 697, 709 Fernández-Pacheco, R. 597 Ferret, P. 125, 139 Ferroni, M. 127, 247, 747 Feuerbacher, M. 459, 645
843
Feuillet, G. 139 Fiawoo, M.F. 117, 205 Fiechter, S. 279 Figge, S. 81 Fischer, R.A. 275 Fitting, H.-J. 17, 529 Flamini, A. 747 Fleurier, R. 117 Flükiger, R. 403 Fonin, M. 621 Fonstad, C.G. 75 Fontcuberta i Morral, A. 295 Fontcuberta, J. 643 Fornara, A. 209 Fourlaris, G. 423, 445, 449, 451 Fournel, F. 665 Foxon, C.T. 47 Frangis, N. 57 Freitag, B. 359 Frenkel, A.I. 281 Fritz, M. 733 Fuchs, D. 593 Fuenzalida, V. 203 Fuess, H. 539 Fujita, Y. 773 Furuya, K. 299, 775
G Gaebler, U. 29 Gaengler, P. 729 Gajović, A. 129 Galerie, A. 415 Galindo, P.L. 45 Galinski, H. 105 Gallo, J. 767 Galtrey, M.J. 41 Galy, J. 543 Gammer, C. 425 Gan, Y. 121 Gao, X.S. 329 Garcia, A. 427 García, M.J. 831 García, R. 69, 75, 77 García-García, A. 607
844
García-González, E. 579 Gass, M. 305, 587, 723 Gassler, N. 727 Gatel, C. 611, 631 Gatica, J.M. 213 Gaudin, G. 613 Gauthier, C. 771 Gautier, E. 33, 613 Gayo, M.D. 805, 831 Geelhaar, L. 83, 143 Geiger, D. 573 Geisler, H. 13 Geist, D. 647 Gemming, S. 573 Gemming, T. 307, 719 Gentile, P. 125, 151 Georgakilas, A. 53, 55 Gerthsen, D. 65, 253, 593, 621 Geserick, J. 219 Ghijsen, J. 141 Gibbs, M.R.J. 599 Gibert, M. 291, 327 Gimel, J.-C. 757 Giorgi, L. 241 Giorgi, M.-L. 333 Giorgio, S. 193 Gloux, F. 89 Gnanavel, T. 599 Godard, O. 755 Godinho, V. 709 Goennenwein, S.T.B. 623 Goetze, F. 29 Goeuriot, D. 529 Golberg, D. 115 Goll, D. 109 Golla-Schindler, U. 265, 809, 811, 825 Golovko, Yu.I. 381 Golubev, Ye.A. 813, 815 Gómez-Herrero, A. 173 González Calbet, J.M. 567, 579 González, D. 69, 77 Gonzalez, L. 91 Gonzalez, Y. 91 Goo, N.H. 109 Gordillo, G. 363
Author Index
Goris, B. 739 Górka, Ł. 453 Gottstein, G. 433 Gouné, M. 477 Goya, G.F. 223 Gradečak, S. 157 Graham, D.M. 41 Gramm, F. 131 Grandjean, N. 79, 637 Grant, D.M. 731 Gregori, G. 565 Gries, K. 733 Griffiths, I. 47 Grill, R. 493 Grobert, N. 169 Gross, R. 623 Grothausmann, R. 279 Grudin, B.N. 471 Gruen, D.M. 167 Grzelczak, M. 153 Gu, L. 243, 659 Guerret-Piécourt, C. 529 Guetaz, L. 207 Gupta, A. 605 Gustafsson, D. 357 Gustafsson, S. 209 Gutmann, E. 569 Gysemans, M. 739
H Habermeier, H.-U. 379 Haeldermans, I. 759 Hagen, C.W. 277 Haghi-Ashtiani, P. 333 Hagiwara, M. 557 Hahn, K. 523 Hamada, E. 511 Hamamoto, C. 763 Han, H. 101 Han, L. 111 Hansen, L.P. 137 Hara, T. 385 Harnchana, V. 601 Hartmann, K. 817
Author Index
Hasanovic, S. 79 Haswell, R. 819 Hauguth-Frank, S. 77 Häusler, I. 71, 83, 133 Häussler, D. 691 He, K. 429, 431 Hébert, C. 531 Hebert, S. 545 Hecq, M. 141 Heeg, T. 325 Heggen, M. 459, 645 Heinrich, W. 817 Heinrichs, J. 741 Helveg, S. 197, 211 Hémono, N. 523 Henderson, G.S. 821 Henry, C.R. 193 Hensel, N. 237 Hermanns-Sachweh, B. 727 Hernández Cruz, D. 753 Hernandez, J.C. 91, 183, 213, 271 Hernández-Velasco, J. 543 Hernando, I. 735 Herrera, M. 69 Herring, R.A. 19 Hervieu, M. 545, 571 Hess, C. 317 Hesse, D. 101, 329 Hessler-Wyser, A. 335, 501 Hetterich, M. 65 Hewitt, I.J. 609 Heyroth, F. 337, 821 Hietschold, M. 365 Hindmarch, A.T. 601 Hirayama, T. 667 Hirmer, M. 95 Hirotsu, Y. 619 Hitchcock, A.P. 753 Hiyama, T. 799 Hjelen, J. 513 Hobbs, L.W. 737 Höche, Th. 523, 821 Hoffmann, M. 693 Hoffmann, M.J. 547 Hofmann, S. 165
845
Holland, M.C. 67 Holmestad, R. 395 Holzapfel, B. 351, 353, 355, 371 Hommel, D. 81 Hondow, N.S. 215 Horak, P. 789 Horibuchi, K. 575 Hörmann, U. 217, 219 Horton, M.A. 321 Hotovy, I. 345 Houben, L. 645 Houdellier, F. 5, 147, 631 Hovmöller, S. 679 Hovsepian, P.Eh. 587 Howe, J.M. 633 Hoyer, I. 729 Hu, W. 433 Huault, T. 303 Hüe, F. 5 Huebner, R. 13 Hug, H. J. 99 Humphreys, C.J. 41, 49 Hünert, D. 689 Hungria, A.B. 183, 213 Hüsing, N. 219 Hwang, S. 245 Hÿtch, M.J. 5, 221 Hyun, Y.-J. 149
I Iacopi, F. 161 Ibarra, A. 223 Ibarra, M.R. 223, 597, 607 Idrissi, H. 649 Ignacova, S. 419 Ikeno, S. 467 Ikuhara, Y. 667 Iliopoulos, E. 55 Ilk, N. 721 Imhoff, D. 103 Immink, G. 159 Infante, I.C. 643 Inkson, B.J. 169, 599
846
Inoke, K. 511 Inoue, K. 463 Ishikawa, T. 763 Iskandar, R. 533 Isshiki, T. 59, 557 Iveland, T. 435 Izgorodin, A. 229
J Jacob, D. 803, 823 Jacques, P.J. 649 Jaffres, P.A. 785 Jäger, W. 109, 669, 691 James, R.D. 413 Jančar, B. 129 Janek, J. 369 Janik, E. 133, 301 Jansen, E. 827 Janßen, A. 825 Janssens, K. 837 Jantou, V. 321 Jensen, S.A. 137 Jentoft, F.C. 237 Jia, C.L. 3, 27, 319 Jia, Y. 535 Jia, Z.H. 435 Jiang, H. 135, 535 Jiang, X. 319 Jimenez, M.C. 709 Jinnai, H. 751, 773 Jin-Phillipp, N.Y. 225 Jinschek, J.R. 761 Johansson, C. 209 Johansson, G.A. 753 Johnson, C.L. 221 Johnson, D.D. 281 Johnson, E. 137, 211, 437 Joly-Pottuz, L. 683 Jordovic, B. 553, 555 Jornsanoh, P. 771 Jouneau, P.H. 25, 139 Jourdain, V. 117 Juhel, M. 37 Juillaguet, S. 57
Author Index
Jurczak, G. 639 Juvé, D. 529
K Kaiser, T. 109 Kaiser, U. 111, 123, 217, 219, 379, 381, 655 Kakas, D. 341 Kalabukhov, A. 297 Kalessaki, E. 651 Kallinen, K. 227 Kamilov, T. 31 Kaneko, J. 453, 455 Kaneko, T. 751 Kanerva, T. 227 Kaplan, W.D. 521 Kappers, M.J. 41, 49 Karakostas, Th. 55, 537, 651 Karczewski, G. 301 Karnthaler, H.P. 385, 425, 465, 481, 489, 505, 647 Karppinen, M. 535 Kasinathan, S. 413 Kašpar, J. 271 Katcho, N.A. 173 Kątcki, J. 61 Katz, H. 229 Kauppinen, E.I. 135, 535 Kawabata, T. 467 Ke, X. 141 Kehagias, Th. 53, 537, 651 Khlobystov, A.N. 113, 123 Khongphetsak, S. 47 Kielbus, A. 439, 441 Kienle, L. 653 Kim, J. 245 Kinnunen, T. 227 Kioseoglou, J. 639, 651 Kirmse, H. 55, 71, 133, 301 Kisielowski, C. 39 Klechkovskaya, V.V. 119, 793 Kleebe, H.-J. 539 Klein, O. 655 Klementová, M. 603
Author Index
Klenov, D. 15 Klimenkov, M. 443 Kling, J. 539 Klingeler, R. 307 Kniep, R. 749 Knote, A. 695 Kobayashi, E. 181 Kobe, S. 617, 627, 713 Kobylko, M. 289 Koch, C.T. 243, 565 Koch, K. 743 Kociak, M. 289 Koguchi, M. 549 Kohout, J. 603 Kokkonidis, P. 445, 449, 451 Kolosov, V.Yu. 343, 657 Komiyama, J. 59 Komninou, Ph. 53, 55, 537, 639, 651 Kong, J.H. 281 Konno, T.J. 447, 619 Konstantinidis, K. 449 Korytov, M. 303 Kosaka, N. 799 Kosiel, K. 61 Kothleitner, G. 387 Kourkoutis, L. 17 Koutsoukis, T. 445, 449, 451 Kovacevic, L. 341 Kovaleva, O.V. 813 Kozhin, A.V. 343 Kraczewski, G. 133 Kralova, D. 263, 765 Krasheninnikov, A.V. 121 Krause, M. 29 Kreiner, G. 407 Kremin, Chr. 693 Kret, S. 133, 301 Krill, G. 103 Krishnan, M. 605 Kröger, R. 733 Kruit, P. 277 Krumeich, F. 541, 579 Kubacka-Traczyk, J. 61 Kübel, C. 733 Kuhn, L.T. 349
847
Kula, A. 453, 455 Kulnitskiy, B.A. 175, 179 Kundu, A.K. 323 Kunert, B. 43 Kungl, H. 547 Kunze, K. 827 Kups, Th. 345, 693, 695 Kuritka, I. 789 Kuskova, A.N. 381
L Lae, L. 417 Lafond, D. 7 Lampke, Th. 685, 687 Lancin, M. 677 Lančok, A. 603 Lançon, F. 665 Landa-Cánovas, A.R. 173, 543 Lange, R. 725 Langenhorst, F. 801 Langlois, C. 187 Lapcikova, M. 767 Lari, L. 143 Łaszcz, A. 61 Lauterbach, S. 539 Le Bouar, Y. 187 Le Guillou, C. 163 Le Pluart, L. 785 Lebedev, O.I. 63, 231, 275 Lebedev, V. 77 Lebius, H. 87 Lee, W. 101 Lefebvre, C. 783 Legendre, F. 397 Legras, L. 411, 427, 483 Leguen, C. 457 Leipner, H.S. 337 Lekston, Z. 491 Lenk, A. 21 Lepistö, T. 227 Lepoittevin, C. 545 Lereah, Y. 269 Leroux, Ch. 233, 347
848
Leroux, F. 739 Leroux, H. 823 Letrouit, A. 571 Leturcq, R. 131 Levin, A.A. 569 Li, J. 753 Li, L. 281 Li, Z.Y. 161, 305 Liang, D. 141 Lichte, H. 21, 269, 363, 569, 629 Licoccia, S. 241 Lim, T. 737 Lima, E. 223 Linck, M. 269, 629 Lindau, R. 443 Lindberg, F. 741 Lipińska-Chwałek, M. 459 Lis, A. 461 Lis, J. 461 Lisiecki, I. 273 Litvinov, D. 65 Liu, Y.-L. 349 Liz-Marzán, L.M. 153, 221, 243, 259 Löffler, D. 253 Löffler, M. 307 Loiseau, A. 117, 187, 205 Lok, M. 249 Lomba, E. 173 Lombardi, F. 357 Longo, P. 67 Loos, J. 769, 795 Lopatin, S. 359 López-Cartes, C. 697 López-Castro, J.D. 199 López-Haro, M. 183, 199, 235 Lotnyk, A. 101 Lozano, J.G. 69, 77 Lu, K. 769 Luca, S. 25 Ludwig, A. 515 Lugstein, A. 149 Lukin, G. 43 Lutsen, L. 759 Luysberg, M. 325
Author Index
M MacFarlane, D.R. 229 MacKenzie, M. 23, 409 MacLaren, I. 547, 581 Mader, W. 191, 623 Madey, T.E. 277 Madigou, V. 233, 347 Maebara, T. 673 Maeda, H. 463 Magén, C. 607 Mahmoud, H.M. 829 Maier, J. 565 Makino, M. 181 Makongo, J.P.A. 407 Malik, S. 609 Malindretos, J. 93 Malo, S. 545 Manca, J. 759 Mangler, C. 425, 465, 481 Manikrishna, K.V. 495 Manolaki, P. 71 Mansouri, S. 87 Mantl, S. 27 Marandian Hagh, N. 525 Marazzi, R. 241 March, K. 103 Maret, M. 11 Marinova, M. 57 Marioara, C.D. 395 Mariolle, D. 25 Marioni, M. 99 Markovich, G. 269 Marquina, C. 597 Marquis, E.A. 473 Marrows, C.H. 601 Martin, D. 637 Martin, J.M. 683 Martínez-Martínez, D. 697 Marty, A. 611, 613 Masenelli-Varlot, K. 25, 771 Massa, W. 799 Massaro, M. 777 Masseboeuf, A. 611, 613 Mateo, A. 479
Author Index
Matlock, D.K. 429, 431 Matos, J.R. 401 Matsuda, K. 467 Matsumoto, H. 447, 773, 775 Matsumoto, K. 283 Matsumoto, T. 549 Mattausch, Hj. 653 Matzeck, Ch. 629 Mayer, G. 621 Mayer, J. 485, 533 Mazeau, K. 781 Mazzucco, S. 289 McAleese, C. 41, 49 McComb, D.W. 23, 177, 321 McFadzean, S. 23 McGilvery, C.M. 23 Méndez Martin, F. 387 Meng, Y. 659 Mensch, K. 819 Menzel, S. 355 Mercey, B. 323 Mercurio, D. 373 Merrifield, R. 305 Meshi, L. 47 Meyer, C. 237 Meyer, D.C. 569 Meyer, J. 39 Mi, S.B. 27 Mickel, C. 355, 371 Miclea, P.T. 551 Midgley, P.A. 91, 183, 723 Miglierini, M. 603 Mihaï, A. 613 Mihailovic, D. 145 Miletic, A. 341 Miljkovic, M. 553, 555 Minkow, A. 217 Minor, A.M. 287 Mira, C. 239 Mirabile Gattia, D. 241 Miron, M. 613 Mishra, R.K. 287 Misják, F. 389 Mitic, V. 553, 555 Mitome, M. 115
849
Mitsuishi, K. 299 Mizera, J. 441 Mizoguchi, T. 667 Mliki, N. 233 Möbus, G. 599 Modin, E.B. 471 Mogilatenko, A. 73 Mohn, E. 257 Molenbroek, A.M. 197 Molina, L. 351, 353, 355 Molina, S.I. 45, 75, 91 Molina-Luna, L. 371 Mompiou, F. 635 Monachon, C. 335 Monnet, I. 87 Monnier, V. 189 Montag, R. 729 Montanari, E. 201 Monteiro, W.A. 401, 421 Montel, J.-M. 835 Monthioux, M. 147 Montone, A. 241 Morales, F.M. 75, 77 Morandi, V. 127, 247 Morante, J.R. 295 Morawiec, M. 469 Morellón, L. 607 Moreno, C. 591 Mori, H. 283 Morimoto, Y. 181 Morin, M. 781 Morselli, L. 837 Moskalewicz, T. 699 Möslang, A. 443 Mouti, A. 79 Muddle, B.C. 399 Muehle, U. 29 Muhammed, M. 209 Mühle, U. 21 Mukhortov, V.M. 381 Müller, E. 131 Muller, K. 723 Muñoz, F. 523 Münzenberg, M. 107 Mur, P. 11
850
Murafa, N. 267 Muralidharan, G. 633 Murray, R.T. 143 Muszalski, J. 61 Muto, S. 575, 577 Mutoro, E. 369
N Najafi, E. 753 Nakamura, J. 467 Nakanishi, H. 59 Nakayama, A. 799 Nasibulin, A.G. 135 Nelayah, J. 243 Nellist, P. 145 Nemeth, I. 43 Neogy, S. 641 Nesper, R. 579 Nespurek, S. 789 Neugebauer, J. 83 Neumann, H.-G. 725 Neumann, W. 55, 71, 73, 133, 301 Newell, D.T. 171 Neykova, N. 765 Nicholls, R.J. 177 Nicolai, T. 757 Nicolosi, V. 145 Nielsen, K. 623 Niemietz, A. 743 Niermann, T. 93 Nietzsche, S. 729 Nikolaidis, K. 707 Nilsen, T. 513 Nishida, I. 577 Nishio, K. 59, 557 Nishioka, H. 763 Nitta, N. 283 Nofz, M. 689 Nolte, P. 225 Nolze, G. 659 Nouet, G. 87 Novak, S. 713 Novikov, S.V. 47 Nowak, C. 261
Author Index
Nozaki, K. 799 Nuzzo, R.J. 281 Nygård, J. 137
O Obergfell, D. 123 Obradors, X. 291, 327, 591 Oehler, F. 125 Ögüt, B. 691 Oh, Y.-J. 245 Ohkura, Y. 763 Oikawa, T. 187, 763 Oksiuta, Z. 503 Olenev, A.V. 63 Olibet, S. 335 Oliver, R.A. 41 Ollivier, A. 333 Olsson, E. 209, 297, 313, 357, 661 Opel, M. 623 Orekhov, A. 31 Ortolani, L. 127, 147, 247 Östberg, G. 661 Otarola, T. 479 Otero-Díaz, L.C. 173 Ozkaya, D. 249
P Pacaud, J. 559 Pailloux, F. 559, 663 Palmer, R.E. 305 Palmquist, A. 741 Pantel, R. 33, 37 Papa, F. 711 Papadopoulou, E. 445, 449, 451 Paraskevopoulos, K.M. 829 Pardo, J.A. 607 Pardoen, T. 509 Parras, M. 567 Pascual, M.J. 523 Pasquini, L. 309 Pastoriza-Santos, I. 221, 259 Pastoriza-Santos, L. 243 Patzke, G.R. 541
Author Index
Pauc, N. 151 Paunovic, V. 553, 555 Pavlidou, E. 55, 701, 703, 705, 707, 829 Pavlovic, V.B. 553, 555 Peiró, F. 295, 643 Peláiz-Barranco, A. 581 Pellicer, E. 311 Peng, Y. 169, 599 Penkalla, H.J. 699 Pennycook, S.J. 45 Perez, M. 457 Perezhogin, I.A. 175, 179 Pérez-Juste, J. 153, 259 Pérez-Munuera, I. 735 Perez-Omil, J.A. 183, 199, 213, 239, 271 Perillat, G. 139 Pesce, E. 777 Peterlechner, M. 489, 505 Petersson, K. 209 Pettersson, H. 357 Pettifor, D.G. 177 Pfund, A. 131 Phiu-on, K. 485 Pichaud, B. 677 Picher, M. 117 Pignot-Paintrand, I. 745 Pileni, M.P. 273 Pillet, J.C. 613 Pintado, J.M. 235 Pireaux, J.J. 141 Piscopiello, E. 241, 309, 777 Plotnikov, V.S. 471 Poelt, P. 779 Pohl, D. 257 Pohl, M.-M. 251 Poissonnet, S. 397 Pokorny, D. 767 Pokrant, S. 9, 51, 195, 683 Polyakov, E.V. 179 Polychroniadis, E.K. 57, 615 Pöml, P. 811 Pongratz, P. 149 Ponzoni, A. 747
851
Popescu, R. 253 Popp, J. 833 Porfyrakis, K. 171 Porter, A.E. 723, 737 Posilović, H. 157 Postava, K. 603 Postigo, P.A. 75 Potapov, P. 13 Powell, A.K. 609 Prellier, W. 323 Presz, A. 301 Pretorius, A. 81 Pritzel, C. 561, 563 Prots, Y. 407 Prusik, K. 469 Pryds, N. 437 Puig, T. 291, 327, 591 Pum, D. 721 Pustovalov, E.V. 471 Putaux, J.L. 781, 783 Putnis, A. 809, 825
Q Quiles, A. 735
R Raanes, M.P. 513 Rabet, L. 499 Radmilovic, V. 367, 473 Radnóczi, G. 389 Rahmati, B. 565 Rainforth, W.M. 587, 711 Ramar, A. 475, 503 Ramírez-Castellanos, J. 567, 805, 831 Ramm, J. 719 Ramos, A.S. 487 Rappaz, M. 501 Raskin, J.P. 509 Ratajczak, J. 61 Rautama, E.-L. 323 Raveau, B. 323 Re, M. 255 Rechenberg, H. 223
852
Recnik, A. 361 Redjaïmia, A. 477, 479 Regula, G. 677 Reibold, M. 569 Reingruber, H. 779 Reiss, G. 107 Reiss, P. 189 Rellinghaus, B. 257, 351, 371 Remmele, T. 83 Renard, K. 649 Renault, O. 151 Rentenberger, C. 425, 465, 481, 647 Retoux, R. 571, 785 Richard, O. 35 Richter, E. 73 Richter, M. 251 Ricolleau, C. 187 Riechert, H. 83, 143 Ripalda, J.M. 45 Rizos, A. 445 Rizzi, A. 93 Rizzo, F.C. 429, 431 Robert, T. 57 Robertson, J. 165 Rodmacq, B. 613 Rodríguez, A. 805 Rodriguez, B.J. 101, 329 Rodríguez-González, J.B. 153, 221, 259 Rodriguez-Manzo, J.A. 121, 155 Rojas, T.C. 709 Romer, S. 99 Rosenauer, A. 81, 733 Rosina, M. 139 Rösler, W. 833 Rösner, H. 359 Ross, C.A. 245 Ross, I.M. 711 Ross, U. 691 Rossell, M.D. 473 Rossinyol, E. 311 Roth, C. 285 Roth, S. 123 Rothe, K. 337 Rother, A. 569, 573, 629
Author Index
Roumeli, E. 829 Rousseau, K. 665 Rouvière, J.L. 33, 85, 125, 665 Rouzaud, J.N. 163 Rozhkova, N.N. 815 Rožman, K.Ž. 617, 627 Ruch, D. 755 Rüdiger, U. 621 Rudolf, C. 671 Rueff, J.M. 785 Ruffenach, S. 69 Rühle, M. 369 Rüssel, C. 523 Russell, B.C. 171 Ruterana, P. 87, 89 Rybacki, E. 827 Ryelandt, L. 649 Ryelandt, S. 649
S Sader, K. 305 Safi, A. 509 Saghi, Z. 599 Sahonta, S.-L. 53, 55 Saijo, H. 361 Saintoyant, L. 483 Sakellari, D. 615 Sales, D.L. 45, 91 Salh, R. 17 Samardžija, Z. 617, 713 Samson, Y. 189 Sánchez, A.M. 45, 91 Sanchez, F. 643 Sanchez, S. 281 Sánchez-López, J.C. 697 Sandino, J. 363 Sandiumenge, F. 291, 327, 591 Sano, H. 773, 775 Santamaría, J. 597 Sarantopoulou, E. 627 Sarro, P.M. 197 Sasaki, T. 575 Sasano, Y. 575 Sato, K. 511, 619
Author Index
Sato, T. 467 Sato, Y. 667 Savalia, R.T. 641 Sberveglieri, G. 127, 747 Schade, M. 337 Schaloske, M.C. 653 Schamm, S. 715 Schaper, A.K. 787, 799 Schappacher, M. 783 Schäublin, R. 475, 503 Schauer, F. 789 Schauer, P. 789 Scherer, T. 217 Schertl, H.-P. 803 Scheu, C. 369 Schils, H. 217 Schletter, H. 365 Schlögl, R. 237, 317 Schmid, H. 191 Schmid, I. 99 Schmidt, B. 17 Schmitt, L. 539 Schmitz, G. 105, 261 Schneider, J.M. 533 Schneider, M. 251 Schneider, R. 65, 253, 621 Schofield, E. 315 Scholz, F. 655 Schöne, J. 669 Schrempel, F. 821 Schröder, F. 275 Schryvers, D. 383, 391, 405, 413, 419, 497, 499, 509, 649, 833 Schubert, J. 325 Schuhmann, H. 93 Schultz, L. 257, 371 Schulze, S. 365 Schwamm, C.L. 343 Schwedt, A. 485 Schweizer, S. 551 Scotchford, C.A. 731 Seeber, B. 403 Seibt, M. 93, 107, 671 Selve, S. 219 Senz, S. 101
853
Serin, V. 631, 717 Servanton, G. 37 Seydoux-Guillaume, A.-M. 835 Sharma, R. 165 Sharp, J. 315 Sheets, W.C. 323 Shibata, N. 667 Shih, S.-J. 519 Shima, T. 557 Shimojo, M. 299, 775 Shiojiri, M. 361 Shorubalko, I. 131 Siemes, H. 827 Sigle, W. 109, 243, 565 Sigumonrong, D.P. 533 Silly, F. 171 Simões, S. 487 Simon, A. 653 Simon, J. 623 Simon, J.P. 11 Simon, P. 749 Simonsen, S.B. 211 Sittner, P. 419 Skepper, J. 723 Skolianos, S. 701, 705 Skoric, B. 341 Slabzhennikov, E.S. 471 Sleytr, U.B. 721 Slouf, M. 263, 765, 767 Smalc-Koziorowska, J. 53 Snauwaert, J. 739 Snoeck, E. 5, 221, 607, 631 Snoek, E. 611 Sobota, J. 453, 455 Soda, M. 95 Sojref, R. 689 Solberg, J.K. 513 Solórzano, G. 203 Soltan, S. 379 Sommer, D. 265 Song, M. 299, 775 Sontakke, P. 605 Sørensen, C.B. 137 Sosna, A. 767
854
Sourmail, T. 457 Sourty, E. 159, 795 Spaldin, N. 573 Speer, J.G. 429, 431 Spiecker, E. 367, 669, 691 Spieß, L. 345, 693, 695 Spiradek-Hahn, K. 493 Spirkoska, D. 295 Srivastava, A.P. 625 Srivastava, D. 625, 641 Srot, V. 109, 369 Stadelmann, P. 79, 637 Steiner, G. 489 Stephan, O. 167, 183, 205 Stergioudis, G. 701, 705, 707 Stierle, A. 225 Stodolna, J. 823 Stöger-Pollach, M. 97 Stolz, W. 43 Stolze, L. 671 Stordeur, M. 337 Störmer, M. 691 Stöver, H. 753 Strand, H. 313 Štrichovanec, P. 607 Strondl, C. 711 Stróż, D. 491 Stuart, M.C.A. 791 Šturm, S. 129, 585, 627 Su, D.S. 237, 317 Šubrt, J. 267 Suenaga, K. 1 Sugamata, M. 453, 455 Sugimori, H. 773 Sugiyama, A. 463 Sukedai, E. 673 Sun, J. 675 Sun, J. 679 Sun, L. 121, 155 Suvorova, E.I. 31, 119, 315, 793 Suzuki, S. 59 Svensson, K. 297, 313 Swinnen, A. 759 Szalay, G. 493
Author Index
Szarpak, A. 745 Szatmáry, L. 267 Szwarcman, D. 269
T Tadano, T. 447 Taguchi, E. 463 Takahashi, Y. 549 Takeguchi, M. 299 Takeuchi, Y. 575 Tambe, M. 157 Tanabe, T. 773 Tanaka, A. 283 Tanaka, M. 299 Tang, D. 769 Tao, T. 731 Tapfer, L. 777 Tarcea, N. 833 Tatsumi, K. 575, 577 Terrones, H. 155 Terrones, M. 121, 155 Tessonnier, J.-P. 317 Tewari, R. 495, 641 Texier, M. 677 Thapa, S.B. 655 Thayne, I.G. 67 Thersleff, T. 351, 355, 371 Thibault, J. 205 Thiel, K. 107 Thiemig, D. 685 Thollet, G. 25, 771 Thomas, A. 107 Thomas, J. 719 Thomas, S.G. 15 Thompson, C.V. 245 Thomsen, P. 741 Tian, H. 383, 497 Tietema, R. 711 Tirry, W. 405, 499 Todaka, Y. 385 Todros, S. 127 Tolley, A. 473 Tonejc, A. 157 Tonejc, A.M. 157
Author Index
Toplišek, T. 713 Torres-Pardo, A. 579 Touzin, M. 529 Trasobares, S. 183, 199, 213, 235, 271 Traversa, E. 241 Tréheux, D. 529 Trettin, R. 561, 563 Trolliard, G. 373, 527 Truche, R. 9 Tsakaloudi, V. 615 Tsiakatouras, G. 53 Tsiaoussis, I. 57 Tsilika, I. 537 Tsuchiya, K. 385, 505 Tsuji, M. 797, 799 Turner, S. 273, 275 Tzormpatzdi, V. 423
U Ubben, K. 107 Uecker, R. 73 Uglietti, D. 403 Ukyo, Y. 575 Umemoto, M. 385 Urai, J.L. 807 Urban, K. 3, 27 Urones-Garrote, E. 173 Utess, D. 13
V Valkeapää, M. 535 van Aken, P.A. 109, 225, 243, 369, 523, 565, 659, 821 van Bavel, S. 795 Van de Casteele, E. 837 Van Den Broek, W. 405 van der Laak, N.K. 41 van Dorp, W.F. 277 van Enckevort, W.J.P. 159 van Esch, J.H. 791 Van Haesendonck, C. 739 Van Humbeeck, J. 497
855
Van Marcke, P. 35 Van Tendeloo, G. 63, 141, 231, 273, 275, 545, 739 Vanderzande, D. 759 Vannod, J. 501 Varela, A. 567 Varela, M. 45 Vargas, J. 223 Veeramani, H. 315 Veleva, L. 503 Velickov, B. 73 Vennéguès, P. 71, 303 Veretennikov, L.M. 343 Verheijen, M.A. 159 Verheyen, P. 15 Viana, F. 487 Vidal, D.M. 213 Vidal, L. 805, 831 Vieira, M.F. 487 Vieira, M.T. 487 Vila, A.L. 613 Vila, E. 543 Villain, S. 233 Villaurrutia, R. 547, 581 Vion-Dury, B. 207 Vippola, M. 227 Vittori Antisari, M. 241, 255, 309, 777 Vlassak, J.J. 509 Vlieg, E. 159 Vlkova, H. 263 Vogel, K. 629 Voitenko, O.V. 471 Volpi, F. 51 Volz, K. 43 Vomiero, A. 127 Vourlias, G. 701, 703, 705, 707 Vovk, A. 607 Vovk, V. 105 Vrejoiu, I. 101, 329
W Waitz, T. 385, 489, 505, 507 Wall, A. 583 Wall, D. 255
856
Walter, M. 107 Walther, T. 375, 377 Walton, M. 737 Wandelt, K. 743 Wang, B. 509 Wang, D. 317 Wang, H. 737 Wang, L.L. 281 Wang, P. 331 Wang, Q. 281 Wang, X. 509 Wang, Z.W. 161 Wantai, Y. 339 Warin, P. 613 Warot-Fonrose, B. 631 Watanabe, M. 369 Weeks, D. 15 Wegscheider, W. 95 Weirich, T. 433 Weiss, P. 253 Welland, M. 723 Wenqing, H. 339 Wepf, R. 131 Wernicke, T. 73 Weyers, M. 73 Weyland, M. 399 Widrig, B. 719 Wieczorek, P. 461 Wiedwald, U. 111 Wielage, B. 685, 687 Wiese, N. 391 Wiesmann, J. 691 Wilcoxon, J.P. 305 Wilde, G. 359 Willinger, M. 317 Wirth, R. 817, 835 Włoch, G. 453, 455 Wojtowicz, T. 133, 301 Wolf, D. 21, 29, 629 Wollgarten, M. 279 Wouters, Y. 415 Wunderlich, R. 217 Wurstbauer, U. 95
Author Index
X Xia, J.H. 319 Xing, H. 675 Xiong, X.-C. 477
Y Yakshinskiy, B. 277 Yamada, K. 511 Yamamoto, K. 181 Yamamoto, T. 667 Yang, J.C. 281 Yang, Z. 833 Yao, L.D. 621 Yasuda, H. 283 Ye, F. 209 Ye, J. 287 Ying, Z. 339 Yokayama, T. 673 Yoshioka, T. 797, 799 Young, N. 519 Young, T.D. 639 Yu, Y.D. 513 Žagar, K. 585
Z Zagonel, L.-F. 151 Zakharov, N.D. 101 Zaleszczyk, W. 301 Zalkind, S. 277 Zanardi, S. 201 Zandbergen, H.W. 197 Zankel, A. 779 Zarnetta, R. 515 Zaspalis, V. 615 Zehl, G. 279 Zeile, U. 819 Zelaya, E. 515 Zhang, D. 679 Zhang, WZ. 659 Zhang, Z. 281, 413
Author Index
Zhang, Z.L. 111, 379 Zhao, Q.T. 27 Zhaoxi 339 Zhigalina, O.M. 381 Zhou, Z. 587 Zhu, D. 41 Zhu, T. 637 Zhu, Y.Q. 293 Zielińska-Lipiec, A. 517 Ziemann, P. 111
857
Zils, S. 285 Zivkovic, Lj. 553, 555 Životský, O. 603 Zolotarevova, E. 767 Zorba, F. 829 Zormalia, S. 451 Zou, X. 679 Zschech, E. 13 Zweck, J. 95, 589, 595 Zysler, R. 223
Subject Index 1 2024 Al alloy 673 2411 371 2-beam conventional TEM 417 3D analysis 837 3D array of chain-folded lamellae structures 775 3D reconstruction 405
A ab initio structure determination 571 aberration corrected TEM 257 aberration correction 221, 331, 619 aberration-corrected electron microscopy 39 absorption 305 acicular ferrite 477 ADF-STEM 395 adhesion 415 AEM 627 α-Fe2O3 293 AFM 79, 87, 617, 693, 713 AFM-TEM 115 a-GaN 637 Ag-Cu alloy 389 ageing 445 air sensitive materials 199 AJ62 441 Al – Si alloys 455 AlGaAs 61 AlGaN 49, 143 AlGaN/GaN DBRs 81 AlInN 79 alloy 79 alloying elements 701 Al-Mg-Ge alloy 395 AlN 59, 87 AlPdMn 635 alteration mechanism 825 alternating copolymer 799
alumina 529, 689, 711 aluminate 571 aluminium alloy 399, 401, 421, 435, 467 aluminium nitride 655 aluminum matrix composite 453 amber 813 amorphization 489, 505 amorphous alloys 471 amorphous carbon 697 amorphous/crystalline interface 107 amorphous-crystalline transformations 657 amylose 781 analysis 461 analytical TEM 301, 565, 719 anatase 219 antiphase boundaries 481 apatite 321, 809 apatite-gelatine 749 artificial pinning centers 371 artist materials 805, 831 atom probe tomography 105, 261 atomic force microscopy 747, 815 atomic layer deposition 715 atomic structure 117, 639 atom-probe-tomography (APT) 473 Au clusters growths 253 austenite- martensite- bainite islands 461 austenitic stainless steel 427
B bacterial cellulose gel-film 119 β-Al3Mg2 517 band gap smearing 87 barium chloride nano-crystals 551 bariumtitanat 583 basal stacking faults 89 BaTiO3 553, 555 B-doped CeO2 565
860
beam damage 811 bend contours 343 β-FeOOH 293 biocompatibility 741 bio-composites 749 biodiesel 227 bio-implant interfaces 741 biomaterial 729, 731, 739 biomineralisation 721, 749 bioresorbable polymers 737 biosensor 127 bismuth ferrite 129 bismuth oxide 129 bismuth-molybdenum oxides 543 block copolymers 751, 783 bond-orientational ordering 787 BOPP 339 bright-field contrast 635 BZO 291
C C 39 C/Cr PVD coating 587 cadmium sulphide 229 Cahn-Hilliard-theory 105 calciumsulfate 561 calciumsulfate-hemihydrate 563 carbide 431, 457, 493, 511 carbide-derived carbon 173 carbon 121 carbon black / polypropylene electron conductive composites 773 carbon materials 123 carbon nanohelices 319 carbon nanomaterials 177 carbon nanoparticles 163, 723, 815 carbon nanostructures 241 carbon nanotube 141, 147, 165, 169, 179, 205, 313, 317 catalysis 211, 233, 251, 317 catalyst 165, 183, 185, 205, 227, 241, 255, 609 cathodoluminescence 17, 361, 789 cation disorder 539
Subject Index
cavitation 755 CdS 39 Ce3+ 577 CeO2 reduction 199 CePrOx Catalysts 235 ceramics 581 ceria mixed oxides 213 CeZrO2 catalysts 271 CGO 327 channeling 681 charge injection 529 chemical 683 chemical composition 467 chemical interaction 141 chemical order 619 chemical solution deposition 327 chemistry 433 CIGS 661 clathrates 531 clay 807 CMOS 7 coated conductors 351, 353, 355 coating 389, 681, 699, 701, 705, 709, 719 coating materials 697 cobalt phases 245 Co-doped ZnO 621 CoFeB/MgO/CoFeB 601 coherence 19 collagen 321 colloids 263 columnar grains 509 complex metallic alloys 459, 517, 645 composites 755 compositional analysis 67, 83 compositional gradient 497 confocal 757 Co-Ni alloy 447 CoNiAl shape memory alloy 391 copper 203 co-precipitation 671 CoPt nanoparticles 187 CoPt thin films 617 core and low loss spectra of TiO2 anatase 195
Subject Index
core shell nanospheres 261 core/shell precipitates 473 core-shell 597 core-shell nanoparticles 305 Cr and U oxides nanoparticles 315 Cr2AlC 533 crack 411 creep 483 cross section polisher 341 cryo-electron microscopy 791 cryo-microscopy 249 cryo-negative staining 783 cryo-TEM 783 crystal defects 73 crystal growth 47 crystal structure 63, 131, 467 crystallization 601 crystallographic boundaries 647 Cs Corrected TEM 147 Cs correction 123 Cs probe corrected TEM 195 Cs-corrected HRTEM 217 Cs-corrector 111, 379 CSD 291 CSL and DSC lattices 479 Cu dispersion 251 Cu-Co alloys 437 Cu-O chain 557 cuprate superconductor 557 Cu-Te amorphous condensates 343 CVD deposition 165
D dark-field diffraction 13 dedicated specimen holders 289 defect 1, 3, 43, 679 deformation 451 deformation defect 675 deformation mechanism 827 deformation twinning 485, 499 degradation 207, 285, 575, 759 dendrites 217 dental fillings 729 dewetting 245
861
DFT 573 diamond 833 diesel engine 211 differential phase microscopy 589 diffraction 525, 559, 681 diffraction contrast 191 diffuse intensity 495 diffuse interfaces 521 diffuse scattering 559 diffusion 151 diffusion couple 501 dilute nitride 43, 669 diopside 537 dislocation 79, 397, 459, 475, 635, 637, 639, 649, 655, 663, 803 dislocation blocking 669 dislocation core 677 dislocation core structure 645 dislocation reduction 47 dislocations structures 427 disordered carbon 173 disordering 63 disordering of ordered structures 481 dissolution-reprecipitation 825 domain texture 323 domain wall 549 domain wall propagation 613 doped-alumina 235 doping contrast 25 double helical structure 751 Dual Beam 427 dual phase 423 duplex (δ+γ) stainless steel 479 duplex stainless steel 659 dynamic recrystallisation 827
E earth analogues 163 EBID 277 EBSD (electron backscatter diffraction) 365, 485, 547, 649, 827 edge dislocation 231 EDS 393, 555, 681, 727, 763 EDS analysis 673
862
EDX (energy-dispersive X-ray spectroscopy) 109, 143, 191, 337, 369, 375, 377, 403, 671, 819 EELS (electron energy loss spectroscopy) 13, 23, 33, 67, 95, 97, 103, 167, 177, 223, 225, 243, 265, 295, 321, 369, 387, 409, 443, 457, 631, 683, 711, 715, 717, 731, 811 EFTEM 49, 51, 181, 223, 231, 243, 587, 597, 621, 623, 731 Egypt 829 electric fields 749 electrocatalyst 207 electroluminescence 229 electrolytic plasma 725 electron beam degradation 789 electron beam nanofabrication 299 electron beam sensitive specimen 761 electron diffraction 117, 135, 333, 425, 535, 539, 583, 653, 781 electron holography 19, 247, 269, 611, 629 electron irradiation 155 electron irradiation damage 663 electron microscopy 213, 401, 421, 435, 471, 473, 695, 703, 707, 713 electron tomography 13, 35, 91, 187, 275, 279, 471, 599, 769 electronic and mechanical investigations 683 electronic excitation 283 electronic properties 51 electrophorese 695 electroplating 685, 687 Elektron 21 439 ELNES 23, 497, 577, 715, 821 elusive structures 807 EMCD 631 emission control 211 energy-filtered TEM 723 environmental SEM 157, 285, 771, 779 environmental TEM 165, 193, 197, 211 epitaxial strain 329
Subject Index
epitaxy 53, 373, 591 EPMA 513 etching 725 exchange bias effect 99 exothermic reaction 487 exsolution 537 extended crystal imperfection 657 extended defects 149 extraterrestrial carbon 163 extreme compression 155
F fatigue 427 Fe3O4 nanoparticles 223 FEBIP 277 Fe-N 477 FePt 257 FePt nanoparticles 189 ferrites 233, 615 ferroelectric materials 101, 381, 527, 547, 549, 579, 581 ferroelectric nanocrystals 269 ferromagnetism 621 FETEM 433 FIB 9, 371, 393, 403, 419, 503, 509, 741, 809, 819 FIB milling 321 FIB preparation 671 FIB/SEM 349, 405, 731 FIB/TEM 835 FIB-cryo-SEM 807 field emission 169 films 381 fin field effect transistor 35 first principles calculation 577 Fischer Tropsch 249 flash memory cell 29 fluorite structure 543 fluorozirconate glass ceramic 551 Frank-Kasper phases 479 fresh-cut 735 fresnoite 821 friction 697 friction-stir welding 673
Subject Index
FSMA 469 FTIR 829 fuel cell 241 fullerene 1, 113, 171, 181
G Ga in Al 393 GaAs (gallium arsenide) 61, 67, 137 GaAs nanowires 157, 295 γ-Al2O3 281 GaMnAs 95 GaN (gallium nitride) 41, 47, 55, 73, 79, 87, 89, 143, 663 gas atomisation 217 GaSb 45 gas-solid interactions 197 gate dielectrics 715 gate oxide 29 Ge 39 gel structure 757 geomaterials 807 geometric phase analysis 11, 691 glass-ceramics 537 GMR sensors 105 goethite 129 gold 259 gold and palladium nanoparticles 263 gold diffusion 125 gold nanoparticles 247 γ-phase 469 γ-polyketone 799 graded materials 683 grain boundary 665, 667, 807, 817 grain boundary width 377 grain refining 455 granular materials 607 ground 819 group theory 479 growth 205 growth kinetics 83 growth mechanism 117, 165, 319 GTL 249
863
H HAADF 223, 295, 303, 331 HAADF-STEM 247, 541, 579, 667, 795 HAADF-STME and HV-TEM tomography / isotactic polypropylene 775 halite 807 hardness 421, 697 heat treatment 493 helical dislocation 673 helicity 135 hematite 129 hemihydrate 561 HEMT 55 heterogeneous catalysis 197, 237, 281 heterojunction 335 heterostructure 103, 295, 323 hexahedral nano-cementites 319 high dielectric constant 715 high manganese austenitic steel 485 high pressure 175 high resolution 519 high resolution electron microscopy 57, 221, 569 high tensile strength steel 511 high-k dielectric 23 high-pressure minerals 801 high-resolution 5 HO2O3 555 hole drilling 599 holography 5, 569, 573 HREM 239, 641, 781 HR-STEM 33, 359 HRTEM (high resolution transmission electron microscopy) 11, 59, 63, 69, 77, 111, 145, 167, 201, 219, 223, 225, 231, 237, 265, 303, 323, 335, 345, 379, 395, 447, 457, 467, 505, 507, 521, 525, 541, 579, 597, 639, 653, 667, 679, 715, 809, 831 HRTEM and X-ray EDS analysis 315 HRTEM simulation 133, 651 hyaluronic acid 745
864
hybrid layered materials 785 hybrid thin film 339 hydrogen storage 309 hydrothermal synthesis 129, 293 hydroxyapatite 731
I IBAD 341 III nitrides 79 ilmenite 825 impact 833 implants 741 impregnated zeolite 251 impurity phases 31 in situ microscopy 289 in situ optical microscopy 561 in situ TEM 287, 657 (In,Ga)N 71 In2O3 39, 345 InAlN 79 InAs nanowires 131 InAs QDs 45 inclusion complexes 781 incommensurate 581, 665 Inconel 738LC 411 indium 55, 161 indium nitride 69 indium oxide 69, 127, 311 In-doped zinc oxide nanorod 191 inert gas condensation 309 InGaAs 65 InGaAsN 83 InGaN 39, 77 InN 87, 93 InP on GaAs 75 in-situ 307 in-situ electron microscopy 121 in-situ SEM 297 in-situ TEM 297, 633 integrated circuits 7 interdiffusion 109 interface 3, 11, 27, 41, 331, 349, 351, 353, 355, 379, 381, 419, 443, 519, 661, 817
Subject Index
interface structure 59, 373, 433, 491 intermetallic 449 internal crystal lattice bending 343 intrinsic electrostatic potential 21 ion implantation 17, 89, 777 ion tracks 87 ionic conductors 543 iron 169 iron antenna 299 iron oxide 545 iron oxide nanostructures 299 irradiation damage 779
L L10 ordering 189 La@C82 177 LACBED 803 LaCoO3 593 Laplace tension 261 laser welding 501 lattice constants 799 layered structures 545 LCMO 643 lead-free ferroelectrics 539 LiMn1.5Ni0.5O4 525 LiNiO2 575 liquid crystals 787 liquid immiscibility gap 437 liquid metal 313 local atomic and electronic structure 577 long-term annealing 439, 441 Lorentz DPC-STEM 549 Lorentz microscopy 391, 589, 595, 629 low carbon steel 493 low hysteresis 413 LSMO films 591 LVSEM 337, 727
M magnesium 309 magnesium alloy 439, 441
Subject Index
magnetic domains 391 magnetic force microscopy 99 magnetic imaging 589, 595 magnetic materials 631 magnetic nanoparticles 597, 609 magnetic properties 615 magnetic shape memory alloys 629 magnetic tunnel junction 601 magnetite 209 magnetostriction 463 maraging 250 445 martensite 383, 385, 423, 469, 507, 605 materials science 717 MBE 55, 75, 295, 303 MCM-41 339 mechanical alloying 453 mechanical properties 287, 423, 445 melon 735 membrane electrode assembly 285 MEMS 197 mesoporous 311 mesoporous material 219 metadislocations 645 metal contact 141 metal gate 23 metal@MOF-5 275 metal-ceramic interfaces 369 metallic glass 641 metallic nanoparticles 243 Mg78.5Pd21.5 407 MgO 357 micelles 783 microanalysis 375, 617 microbial reduction 315 microcapsules 745 microelectromechanical systemts 197 microflakes 833 microscopies 101 microstructure 27, 383, 423, 429, 437, 439, 441, 445, 447, 463, 553, 615, 685, 687, 689, 709, 735 micro-wire 497 mineral bridge 733 mineralization 737
865
minerals 835 misfit 417 misfit dislocation 359, 651 Mn partitioning 461 Mn4Si7/Si films 31 MOD 591 modeling 781 modulated structure 57, 535 Moiré 371 monochromator 195 morphology 159, 795 morphology of wear particles 767 MOSFET 67 Mössbauer spectroscopy 603 MOVPE 43 MRAM 107 multicompartment nanostructures 791 multiferroics 329, 573 multilayer 109, 487, 691, 699 multi-metal silicides 671 multiple scattering calculation 103
N nacre 733 nano-beam diffraction 15 nanobelts 115 nanobridges 599 nanocasting 311 nanoclusters 245 nano-columns 93 nanocomposite 697, 709, 765, 785 nanocrystal 197, 229, 273, 625, 641 nanocrystalline materials 425 nanocrystalline microstructure 491 nanocrystallinity 465, 489, 719 nanocrystallization 89, 523 nanodiamond 167 nanodots 87 nanofabrication 599 nano-filaments 117 nanohuts 327 nanointerface 341 nanomagnetism 627 nanometer scale 277
866
nanometer size wear debris 767 nanoparticle 111, 203, 221, 225, 231, 233, 257, 265, 267, 281, 283, 603, 619, 685, 687, 739, 833 nanoparticle arrays 721 nanoparticle lattice parameter 239 nanorod 139, 259, 293 nanospheres 627 nanostructures 101, 255, 349, 585, 595, 725 nanotips 599 nanotube 1, 113, 115, 117, 121, 135, 307, 753 nanotube growth 155 nano-twinning 593 nanowhiskers 219 nanowire 115, 133, 137, 139, 143, 145, 149, 151, 159, 167, 301 NBT 527 NEXAFS 753 Ni clusters 255 Ni2MnGa 463 Ni4Ti3 405 nickel 153 nickel aluminides 487 Ni-Mn-Ga 605 NiTi 489, 505, 507 NiTiHf 489 nitride 55, 85, 457 nitride interfaces 651 nitride precipitates 387 nitride semiconductors 87 NMR 603 non-conductive samples 529 nonpolar Gallium Nitride 53 non-stoichiometry 567 nucleation 357, 633 nuclei 205
O octahedral tilting 539 ODS steels 443 off-axis electron holography 9 off-stoichiometric alloy 463
Subject Index
oligonucleotide 747 one-pot synthesis 271 onion 175 optical properties 87, 97 order 521 order/disorder phenomena 187 ordered solid solution 389 organic solar cells 759 orientation contrast 827 orientation relationship 479, 659 orthopaedic implants 737 oxidation 415, 703 oxide 3, 475 oxide electrode 347 oxide layer (or oxidation) 531
P pack cementation 701 particles 695 PE 339 peak broadening 425 peapods 147 pearlite 477 PEMFC 207, 285 perovskite 323, 567, 593 perovskite-solid solution 579 perpendicular magnetic anisotropy 611, 613 phase identification 681 phase mapping 69 phase segregation 591 phase separation 389, 523 phase transformation 479, 495 phase transition 527, 583, 787 phason walls 635 photoactive layer 795 photocatalyst 215, 267 piezoresponse force microscopy 269 pigment 819, 829 plan view 239 plasma-jet deposition 603 plasmon loss electrons 19 plastic deformation 459 PLD 371
Subject Index
PMPSi 789 p-n junction 21 poisoning 227 polarity 47, 85 poly(butylene terephthalate) (PBT) 797 poly[methyl(phenyl)silylene] 789 polyamide 765 polycarbonate 777 polyethylene 755 polymer crystallization 797 polymer nanocomposites 769 polymer stabilised nanoparticles 119 polymers 753 porosity 807 potential mapping 565 powder metallurgy 503 precession electron diffraction 803 precipitate 391, 397, 409, 417, 435, 467 precipitation 383, 395, 449, 451, 457, 673 precipitation crystallography 659 prismatic stacking faults 89 probe Cs corrector 33, 359, 665 projected potential 761 protection layers 689 protein assembly 721 Pt 281 Pt catalysts 193 pulsed laser deposition (PLD) 347, 607, 699 PVD 345 pyroxene 823
Q quantification 387, 769 quantitative 757 quantitative electron microscopy 729 quantitative image analysis 773 quantum cascade laser 61 quantum dot 65, 85, 91, 303 quantum rings 91 quantum wells 41, 71
867
quasi in-situ structure research 285 quasicrystals 635 quenching and partitioning 429
R radiation damage 37, 39, 835 Raney-type Ni 217 rapid solidification (RS) 455 rare earth doping 89, 553 rare earth oxides 715 reactive diffusion 261 recrystallization 401, 483 redox process 793 replacement reaction 809 replica TEM preparation 333 retardation 97 RHEED 731 rhodium 225 risk assessment 181 rotational twinning 407 Rubens 805, 831 Ruddlesden Popper phases 569 ruthenium catalysts 213, 279 rutile 219
S Samson phase 517 sandstone characterization 837 sapphire 655 scanning probe microscopy 813 Se/polymer particles 793 segregation 623, 667 selective oxidation 333 selenium nanowires 119 self assembly 273, 743, 791 SEM (scanning electron microscopy) 25, 79, 201, 265, 313, 365, 501, 503, 513, 555, 693, 701, 705, 727, 745, 771, 805, 819, 829 semiconductor silicon 33 semiconductors 5, 9, 25, 37 severe plastic deformation 465, 481, 647
868
shape memory 605 shape memory alloys 413, 419, 515 shape memory effects 491 shear band 385, 465 shells 153 short range ordered cubic phases 567 shungite 815 Si As/P doped 37 Si nanowire 161 SiC (silicon carbide) 15, 57, 677 SiC buffer layer 59 SiC fibers 713 SiGe 15 silica 415, 529 silica layers 17 silicon 37, 151, 747 silicon nanowires 125 silicon oxide 453 silicon polymers 789 simulation 495 single atom detection 39 single crystals 781 slicing view 217 SMA 501 soft materials 763, 779 softmagnetic material 625 sol-gel electrophoretic deposition 585 sol-gel synthesis 219 solid bitumens 813 solid-liquid interface 633 solution treatment 485 soot oxidation 211 specimen preparation 393 specimen surface 21 specimen tilt 375 spectrum imaging 409 spherulite 797 spin torque 107 spintronic 103 Sr4Ru2O9 347 SrTiO3 171, 325, 357, 585 stacking faults 637, 677 stacking sequence 557 stainless steel 415, 501
Subject Index
Stardust 823 steel 409, 429, 431, 457 STEM (scanning transmission electron microscopy) 49, 67, 71, 81, 85, 125, 145, 199, 215, 223, 235, 265, 271, 305, 349, 433, 453, 455, 665, 763, 769, 771 STEM and EDS 793 STEM DF 809 STEM EELS/EDX 37 STEM SEM 785 STEM/EELS 575, 587 STEM-simulation 81 STM (scanning tunnelling microscopy) 171, 297, 307 STM-TEM 115 STO 643 strain 5, 15, 77, 221, 325 strain analysis 13 strain contrast 397 strain engineering 669 stress 547, 615 stress-strain response 649 structural 683 structural change 283 structural properties 363 structure 161, 217, 237 structure determination 541 structured surfaces 743 STXM 753 SU-8 693 sub-nano analysis 511 sulfated zirconia 237 superalloy 411 superaustenitic stainless steel 451 superferritic stainless steel 449 superlattice reflections 539 superlattices 101 surface 725 surface composition 183 surface layers 497 surface Ostwald ripening 253 surface oxidation 225 surface plasmon mapping 243
Subject Index
surface precipitates 659 surface restructuring 193 surface structure 183 swift ions irradiation 87
T TEM (transmission electron microscopy) 27, 45, 53, 61, 65, 67, 75, 79, 81, 87, 89, 113, 123, 131, 137, 153, 175, 179, 185, 189, 209, 215, 227, 229, 253, 259, 263, 273, 275, 295, 317, 329, 363, 367, 371, 399, 403, 415, 427, 439, 453, 455, 463, 489, 497, 499, 509, 513, 515, 523, 527, 531, 533, 539, 545, 571, 585, 603, 607, 621, 625, 643, 649, 661, 691, 699, 733, 741, 745, 759, 761, 777, 785, 801, 819, 823 TEM analysis 149 TEM and X-ray EDS analysis 31 TEM study 407 TEM tomography 795 TEM/SAED 157 tempering 431 TEM-STEM imaging 51 tensile tests 771 texture 365 TG Measurements 703 thermal analysis 401 thermal reaction 105 thermal spray 705 thermal treatments 421 thermoelectric film 337 thermogavimetric measurements 705 thin film 73, 335, 351, 353, 355, 357, 367, 373, 509, 559, 603, 623 thin film characterization 515 thin film diffusion 817 thin films solar cells 363 threshold photoemission 151 Ti6Al4V 499 TiAlSiN 709 TiAlYN/CrN 711
869
TiC 697 tin oxide 127 Ti-Ni-P 681 Ti-Ni-Pd 515 TiO2 765 titanate nanotubes 765 titania 219 titania surfaces 741 titanium 725 titanium alloy 675, 699 titanium dioxide 267 titanoniobates 785 tomography 29, 159, 367, 753 tool steels 703 toxicity 723 transition elements in silicon 671 transition insulator-metal 167 transmission electron microtomography 773 transmission electron tomography 751 transport measurements 289 transport of intensity equation 611 TRIP 485 tungsten 503, 707 twin 203, 385, 803 twinning 413, 479 twinning mechanism 649 TWIP 485 TWIP steels 649
U UHP metamorphism 801 ultrahigh molecular weight polyethylene 767 ultra-low-k 51 UNCD 167 uncompensated spins 99
V vacancies 231 vacancy clusters 399 valence states 811
870
Van Gogh 819 vanadium oxide 317 vapour-liquid-solid growth 133 varistor 361 VEELS 51
W wax crystallisation 743 WDXS 617 wetting layer 377 wires 93 wollastonite 537
X XANES 821 XEDS 515 xidation 705 XPEEM 151 X-ray 375 X-ray absorption 753 X-ray grazing incidence diffraction 11 X-ray micro-CT 837
Subject Index
X-ray optics 691 XRD (X-ray diffraction) 571, 653, 707 XRPD 201
Y Yb2O3 555 YBa2Cu3O7-δ 357 YBaCo4O7+δ 535 YBCO 291, 371 Y-junction 179 yttria 475
Z Z-contrast 665 zeolite 201 zeolite beta 679 zinc 707 zirconium 483 zirconolite 811 (Zn,Mn)Te 301 ZnO 139, 361, 667 Zr3Al 647
