TRIENNIAL REPORT CNR - INFM

Transcription

1 Laboratorio Nazionale MDM Materials and Devices for Microelectronics TRIENNIAL REPORT CNR - INFM

2 present Triennial Report describes the research activities carried out at the MDM National Laboratory, also in collaboration with other academic and industrial institutions, TThe during the years The Report has been divided into eight chapters dealing with the most representative active research lines providing an introduction followed by a synthetic description of selected studies. These are related to the investigation of materials and processes for ultra-scaled logic and memory devices based on the CMOS technology, novel non-volatile memory devices, spintronics, neuroelectronics, and metamaterials. The theoretical effort, in support of the main experimental activities is also reported. All these activities were enabled by the financial support from INFM, CNR, STMicroelectronics, the European Commission, the Ministry of University and Research, and the Ministry of Foreign Affairs. A description of the personnel, the facilities, the main projects, and of the financial support completes the report. Marco Fanciulli Director 1

3 CONTENTS ø Introduction 8 ø Facilities 10 ø Scientific Board 12 ø Executive Board 13 ø Personnel 14 ø Publications 22 ø Conferences 30 ø Organization of conferences 37 ø Research Projects 38 ø International Collaborations 48 ø National Collaborations 49 2

4 1. Scaling issues in CMOS logic and memory devices ø 1 Scaling issues in CMOS logic and memory devices 50 ø 1.1 Structural and electrical properties of Hf-based oxides 52 or interpoly applications C. Wiemer, R. Piagge, S. Spiga, E. Bonera, M. Fanciulli, M. Alessandri, G. Ghidini, M. Caniatti, A. Sebastiani, D. Caputo ø 1.2 Al 2 as interpoly dielectrics in non-volatile memory devices 54 C. Wiemer, R. Piagge, S. Spiga, E. Bonera, V. Fiorentini, M. Fanciulli, M. Alessandri, G. Ghidini, A. Del Vitto, M. Caniatti, A. Sebastiani, S. Alberici, G. Pavia ø 1.3 HfO 2 as gate dielectrics for ultra-scaled CMOS devices 56 S. Spiga, C. Wiemer, G. Scarel, S. Baldovino, G. Tallarida, M. Fanciulli, C. M. Compagnoni, A. S. Spinelli, A. Bianchini, A. L. Lacaita ø 1.4 Nanoscale electrical properties of HfO 2 and ZrO 2 thin films 59 studied by conducting atomic-force microscopy S. Kremmer, H. Wurmbauer, C. Teichert, G. Tallarida, S. Spiga, C. Wiemer, and M. Fanciulli ø 1.5 Thermal stability of HfO 2 /TiN gate stacks for 45 nm CMOS devices 62 C. Wiemer, M. Perego, M. Fanciulli, V. Cosnier, P. Besson, V. Loup, L. Vandroux, S. Minoret, M. Cassé, X. Garros, J-M. Pedini, S. Lhostis, K. Dabertrand, C. Morin ø 1.6 Oxygen diffusion in HfO 2 /SiO 2 /Si stacks 64 S. Ferrari and M. Fanciulli ø 1.7 Characterization of the mechanical stress induced 65 in silicon during device fabrication E. Bonera, M. Fanciulli, G. Carnevale, M. Mariani ø 1.8 Low-k materials for intra-metal dielectrics 67 A. M. Ferretti, C. Wiemer, E. Bonera, C. Rossi, M. Fanciulli ø 1.9 Advanced materials for interconnects 70 G. Tallarida, C. Wiemer, L. Aina, S. Alberici, E. Ravizza, A. Giussani, G. Pavia, E. Varesi, G. Brunoldi, S. Guerrieri, S. Grasso, E. Ravizza, S. Spadoni 3

5 2. Emerging Materials for nanoscale CMOS devices ø 2 Emerging Materials for nanoscale CMOS devices 74 ø 2.1 Rare earth oxides on Si for logic and memory applications: 76 binary and ternary compounds G. Scarel, C. Wiemer, S. Spiga, E. Bonera, G. Seguini, G. Tallarida, X. Li, M. Fanciulli, Y. Lebedinskii, A. Zenkevich, I.L. Fedushkin, V. Fiorentini, F. Boscherini, S.D. Elliott, S. Schamm, G. Pavia ø 2.2 Band alignment at the La 2 Hf 2 O 7 /(001)Si interface 78 G. Seguini, S. Spiga, E. Bonera, M. Fanciulli, A. Reyes Huamantinco, C.J. Först, C.R. Ashman, P.E. Blöchl, A. Dimoulas, G. Mavrou ø 2.3 Dielectrics for channel materials alternative to Si: towards strained-si, 80 Ge and GaAs devices S. Spiga, C. Wiemer, G. Scarel, G. Seguini, M. Perego, S. Ferrari, S. Baldovino, M. Fanciulli, A. Zenkevich, Y. Lebedinskii, Y. Panayiotatos, A. Dimoulas ø 2.4 GeO 2 films grown on Ge subtrates by Atomic Layer 84 Deposition and Molecular Beam Epitaxy A. Molle, M. N. K. Bhuiyan, G. Tallarida, M. Perego, G. Scarel, M. Fanciulli, I. L. Fedushkin, A. A. Skatova ø 2.5 HfO 2 as gate dielectrics for Ge-based devices 86 S. Spiga, C. Wiemer, G. Scarel, G. Tallarida, G. Seguini, M. Perego, S. Ferrari, M. Fanciulli, G. Mavrou, A. Dimoulas, S. Kremmer, C. Teichert, G. Pavia ø 2.6 Epitaxial HfO 2 on high-mobility semiconductors: theory and experiment 89 C. Wiemer, A. Debernardi, G. Scarel, M. Perego, M. Fanciulli ø 2.7 Epitaxial Gd 2 films on Ge 92 A. Molle, M. N. K. Bhuiyan, G. Tallarida, C. Wiemer, M. Fanciulli, G. Pavia 3. Innovative Technologies in Non-Volatile Memories ø 3 Innovative Technologies in Non-Volatile Memories 94 ø 3.1 Si nanocrystals in a dielectric matrix: synthesis and characterization 96 M. Perego, S. Spiga, M. Fanciulli, C. Bonafos, and G. Benassayag 4

6 ø 3.2 Metallic nanocrystals embedded in SiO 2 : low temperature CEMS 98 characterization R. Mantovan, S. Spiga, A. Debernardi, and M. Fanciulli ø 3.3 Structural and functional properties of phase change 100 materials grown by sputtering and MOCVD C. Wiemer, M. Fanciulli, C. Giessen, R. Bez, A. Pirovano, S. Rushworth, J. Siegel, C. N. Afonso, A. Abrutis, V. Plausinaitiene ø 3.4 Thermal characterization on phase change materials 102 J. L. Battaglia, A. Teren, C. Monguzzi, E. Varesi, R. Cecchini, C. Wiemer, R. Fallica, A. Kusiak, C. Rossignol, N. Chigarev, S. Cocco ø 3.5 Binary oxides for resistive switching non volatile memories 105 S. Spiga, C. Wiemer, G. Scarel, M. Perego, G. Tallarida, S. Ferrari, A. Cappella, H. Lu, M. Fanciulli ø 3.6 Electrical and chemical investigations of ZnO layers grown by Atomic 107 Layer Deposition for selectors in crossbar non volatile memories S. Ferrari, E. Speets, N. Huby, A. Pirovano, E. Guziewicz, A. Wójcik, M. Pra, P. Lugli 4. Spintronics ø 4 Spintronics 110 ø 4.1 Atomic Layer Deposition for magnetic tunnel junctions 112 M. Georgieva, H. Lu, M. Perego, G. Scarel, A. Zenkevich, M. Fanciulli ø 4.2 Characterization of Fe/high-κ oxide interfaces 114 R. Mantovan, C. Wiemer, A. Zenkevich, M. Fanciulli ø 4.3 Shallow donor electron spin coherence and manipulation in Si and SiGe 117 M. Fanciulli, A. Ponti, A. Ferretti ø 4.4 Study of the static magnetic field and microwave irradiation response 119 of the random telegraph signal in MOSFETs for qubit implementation E. Prati, M. Fanciulli, G. Ferrari, M. Sampietro ø 4.5 On-line and off-line Mössbauer spectroscopy investigation 122 of the magnetic properties of oxides R. Mantovan, M. Fanciulli, G. Weyer, H.P. Gunnlaugsson, D. Naidoo, R. Sielemann, K. Bharuth-Ram, T. Agne 5

7 5. Devices based on organic and polymeric semiconductors ø 5 Devices based on organic and polymeric semiconductors 124 ø 5.1 Al 2 gate dielectric in poly (3-hexylthiophene) based transistors 126 E. Peron, F. Perissinotti, G. Tallarida, S. Ferrari, L. Fumagalli, D. Natali, and M. Sampietro ø 5.2 Atomic Layer Deposited Al 2 as a capping layer 127 for polymer based transistors S. Ferrari, F. Perissinotti, E. Peron, L. Fumagalli, D. Natali, and M. Sampietro ø 5.3 A novel method for the production of top contact thin film 128 transistors based on organic semiconductors S. Ferrari, F. Perissinotti, E. Peron, L. Fumagalli, D. Natali, and M. Sampietro 6. Very high-κ oxide for Neuroelectronics ø 6 Very high-κ oxide for Neuroelectronics 130 ø 6.1 Electrical properties of neuroelectronic devices with various interlayers 132 between TiO 2 and p-si(100) F. Wallrapp, G. Scarel, M. Perego, G. Seguini, M. Fanciulli, and P. Fromherz ø 6.2 Energy band alignment at the TiO 2 /Si interface 133 with various interlayers M. Perego, G. Seguini, G. Scarel, M. Fanciulli, F. Wallrapp, and P. Fromherz 6

8 7. Metamaterials ø 7 Metamaterials 136 ø 7.1 Theory of propagation in negative refractive index media 138 E. Prati ø 7.2 Experiments on transmission of negative refractive index media 140 based on Split Ring Resonators C. Amabile and E. Prati 8. Tailoring innovative devices by parameters-free simulations ø 8 Tailoring innovative devices by parameters-free simulations 142 ø 8.1 Semiconducting and high-κ oxides for ultra-scaled devices and spintronics 144 A. Debernardi and M. Fanciulli ø 8.2 The magnetic map of Mn-based thin film alloys on Ni substrates 146 B.R. Malonda-Boungou, B. M Passi-Mabiala, A. Debernardi, S. Meza-Aguilar, C. Demangeat ø 8.3 Parameter free calculation of shallow states in external field 147 A. Debernardi, M. Fanciulli, and A. Baldereschi ø 8.4 Heterojunction for spintronic devices 149 A. Debernardi, M. Peressi, and A. Baldereschi 7

9 Introduction TThe MDM (Materials and Devices for Microelectronics) Laboratory is an Italian National Laboratory belonging to INFM, the Italian Institute for the Physics of Matter since 2006 part of CNR, the National Research Council. MDM is a state-of-the-art facility located within the STMicroelectronics complex at Agrate Brianza, near Milan (Italy), and its major research activities are focused on investigating the structural, electrical, optical and magnetic properties of materials for present and future nanoelectronics, as well as developing innovative processes and characterization techniques. The cooperation between MDM and STM, one of the most important semiconductor industries worldwide, began in 1996 with the establishment of a small laboratory at Agrate, which is one of STM s largest multi-functional sites, hosting manufacturing, R&D (Research and Development), product design and marketing activities. The importance of the research performed at MDM derives from the most critical trend in electronics, namely the continually decreasing dimensions of the transistors and other components used to build microand nano-electronic devices, the System-on-Chip devices that integrate tens or hundreds of millions of transistors on a tiny silicon chip and provide the heart of electronic applications such as mobile phones, settop boxes and car engine management units. Typically, these dimensions are reduced by around 30% every two years. The reason why the global semiconductor industry pursues this goal of continual reduction is that reducing the size of the transistors used to build complex devices makes the resulting devices faster, cheaper and less power-consuming, thereby enabling the expansion of existing markets and the creation of new applications. However, the development of a new generation of semiconductor technology often involves the introduction of specialized new materials and it is essential that the properties of these materials are thoroughly understood and characterized before they can be introduced into a high volume manufacturing flow that is both complex and highly cost-sensitive. 8

10 The strong collaboration with the industrial partner should also be seen as an opportunity for scientists to address fundamental issues in condensed matter, made possible by the availability of advanced and state of the art industrial devices and prototypes. The research activity carried out at the MDM Laboratory is a delicate balance among: ø long-term actions aimed at the development of materials, processes, as well as advanced characterization techniques and methodologies for future scenarios in nano-electronics; ø medium-term actions strongly connected with issues in these fields and with potential impact on the next generation devices; ø short-term actions aimed at solving problems encountered by the industrial partner in its R&D and production activity. With the national and the international collaborations established during the years the capabilities of addressing different problems, relevant for industrial applications, are strongly enhanced and, at the same time, new fundamental research activities are stimulated. The Laboratory is also strongly active in training students and young researchers, at the national and international level, and is involved in several seminars, schools, and courses. The training activity focuses on fundamental as well as more applied and industrial-related issues. Mainly due to the external projects started in the last three years the personnel including staff researchers and technicians, post-doctoral fellows, Ph.D. students, undergraduate students, and administrative personnel, grew from 18 persons to 40. 9

11 Facilities facilities comprise 360 TThe m2 of laboratories, including a 94 m 2 class 1000 clean room, plus administration and support offices. Further extensions are planned for Advanced Materials Growth and Process Facilities ø 94 m 2 clean room class 1000 (Extension to 120 m 2 in 2007) ø Atomic layer deposition (ALD) with line (4 wafers) ø Atomic layer deposition (ALD) with line (8 wafers) ø Molecular beam epitaxy (MBE) equipped with O source ø Metallorganic chemical vapour deposition (MOCVD) (Feb.2007) ø Cluster tool (ALD, evaporator, sputtering, preparation) (2007) ø Photo-lithography ø Electron-beam lithography ø Thermal and e-beam evaporator ø Rapid thermal processing (RTP) and furnace annealing systems ø Wet bench for wafers cleaning ø X-ray irradiation

12 Characterization Facilities ø Scanning electron microscope (SEM) ø Scanning probes in air or controlled atmosphere: - AFM, STM, MFM/EFM, SCFM, KPFM - SThM - SNOM ø Scanning probes in UHV and at variable temperature ( K): - AFM, STM, KPFM, BEEM ø X-ray diffraction (XRD) and reflectivity (XRR) ø Total reflection X-ray fluorescence (TXRF) and reflectivity (XRR) (Oct. 2006) ø X-ray photoelectron spectroscopy (XPS) ø Low-energy ion scattering spectroscopy (LEIS) ø ToF-SIMS (Access to the STMicroelectronics system) ø Electrical characterization ( K): - I-V, C-V, G-V - Deep-level transient spectroscopy (DLTS) and Laplace-DLTS - Noise - Internal photoemission spectroscopy (IPE) - Inelastic electron tunnelling spectroscopy (IETS) - Hall effect (2007) ø Optical Spectroscopies ( K): - Micro-Raman with excitation in the visible (488 nm, 633 nm) - Micro-Raman with excitation in the UV (347 nm) - Photoluminescence (PL) spectroscopy - Fourier-transform infrared spectroscopy (FTIR) (middle- and far-ir) - Spectroscopic ellipsomtery (SE) ø Electron spin resonance spectroscopy ( K): - X-band and Q-band CW-ESR - Pulse EPR in X-band (FT-EPR) - Electrically detected magnetic resonance (EDMR) - Electron-nuclear double resonance (ENDOR) in X band - Multi-frequency EDMR (up to 40 GHz, upgrade up to 200 GHz in 2007, fields up to 12 T, temperatures down to 260 mk) ø 57 Fe and 119 Sn conversion electron Mössbauer spectroscopy (CEMS) (120 K and room temperature set-ups) Computational Facilities: ø 6 nodes of a 32 node cluster E4-InfiniNode (2 CPU AMD Opteron 250, 4 GB di RAM) located at the National supercomputer facility CASPUR

13 SCIENTIFIC BOARD (Since 2007) ø Paolo Giuseppe Cappelletti FTM Vice President Non Volatile Memories Technology Development STMicroelectronics Srl Via C. Olivetti Agrate Brianza (Mi), Italy Phone: Fax: [email protected] ø Alain Claverie Directeur de Recherche CNRS CEMES-CNRS Groupe Nanomatériaux 29, rue Jeanne Marvig, BP Toulouse Cedex 4, France Phone: Fax: ø Paolo Lugli Professor Lehrstuhl für Nanoelektronik Technische Universität München Arcisstrasse 21 D München, Germany Phone Fax: ø Theodore M. Moustakas Professor of Electrical and Computer Engineering Professor of Physics Director of Wide Bandgap Semiconductor Laboratory Boston University 8 St Mary s St Boston, MA 02215, USA Phone: Fax: [email protected] [email protected] [email protected] ø Michael Pepper Head of the Semiconductor Physics Group Cavendish Laboratory J.J. Thomson Avenue, Madingley Road, Cambridge CB3 0HE, United Kingdom Phone: +44 (0) Fax: +44 (0) Secretary: +44 (0) [email protected] ø Matthias Wuttig [email protected] Professor Physikalisches Institut IA RWTH-Aachen D Aachen, Germany Phone: Fax

14 EXECUTIVE BOARD (Since 2007) ø Marco Fanciulli Director CNR-INFM MDM National Laboratory Via C. Olivetti Agrate Brianza (Mi), Italy Phone: Fax: ø Sabina Spiga [email protected] Researcher CNR-INFM MDM National Laboratory Via C. Olivetti Agrate Brianza (Mi), Italy Phone: Fax: ø Grazia Tallarida [email protected] Researcher CNR-INFM MDM National Laboratory Via C. Olivetti Agrate Brianza (Mi), Italy Phone: Fax: ø Giorgio De Santi [email protected] FTM - Director Agrate R2 Operations and Manufacturing Process Development STMicroelectronics Srl Via C. Olivetti Agrate Brianza (Mi), Italy Phone: Fax: ø Pietro Palella General Manager STMicroelectronics Srl Via C. Olivetti Agrate Brianza (Mi), Italy phone: Fax: ø Alberto Modelli Physics & Materials Characterization Manager STMicroelectronics Srl Via C. Olivetti, Agrate Brianza (Mi), Italy Phone: Fax: [email protected] e.mail: [email protected] 13

15 Personnel Marco Fanciulli, Director of Research Marco Fanciulli graduated cum Laude in Nuclear Engineering at the Politecnico of Torino (Italy) in 1987 and obtained his PhD in Applied Physics from Boston University, Boston (USA) in 1993 for his work on the characterization of wide-band-gap semiconductors using magnetic resonance spectroscopies. In 1993 he joined the Institute of Physics and Astronomy at the University of Aarhus (Denmark) as an assistant professor to conduct research on defects in silicon using different techniques (DLTS, CEMS, EPR) and on transition metal silicides. In 1997 he was appointed associate professor. From 1998 he is the director of the MDM Laboratory. He has published more than 150 papers on the growth, by MBE, PLD, ALD, and characterization, using different spectroscopy methods most of them related to hyperfine interaction and magnetic resonance detection, of different materials. His current interests are related to the characterization of high-κ materials and of nanocrystals embedded in silicon oxide, to the development of advanced magnetic resonance characterization techniques, to the development of silicon based qubits, and to other fundamental issues related to materials and devices for nanoelectronics and spintronics. He has been the coordinator of the European FET assessment project ESRQC (Electron Spin Resonance for Quantum Computing) and of industrial projects with STMicroelectronics. He has been the principal investigator in two EC projects on nanocrystals (GROWTH: NEON Nanoparticles for electronics) and on epitaxial oxides (IST: INVEST Integration of very high-κ dielectrics with silicon CMOS technology) as well as in other national projects. Currently he is the principal investigator in three EC projects: ET4US Epitaxial technologies for ultimate scaling, REALISE Rare earth oxide atomic layer deposition for innovation in electronics; and EMMA Emerging Materials for Mass-storage Architectures. In the recent years he organized several international workshops and symposia. Marco Fanciulli has been also professor of semiconductor physics at the Università degli Studi of Milano (Italy) since Alberto Debernardi, Senior Researcher Alberto Debernardi has mainly developed his professional experience within the framework of density functional theory. In 1995 he received his PhD at the SISSA - Trieste (Italy) presenting a new method to study the anharmonic effects in crystals from density functional perturbation theory; he applied his method to compute the phonon lifetime of semiconductors. After a postodctoral experience at the MPI-Stuttgart were he collaborated with Cardona and Parrinello groups on vibrational properties of semiconductors, and at the IPCMS of Strasbourg were he computed ab initio magneto optical and thermodynamics properties of metals, he was researcher ( ) at INFM, at Trieste University/SISSA where he studied by first principles magnetic semiconductors, interfaces and magnetic heterostructures. In 2001, he took the Degree Habilitation à Diriger des Recherches, at the Strasbourg University. During the July 2002 and 2003 he was invited professor at the Tours University. Since summer 2004, he is senior researcher at MDM Laboratory, his present research interests include high dielectric constant oxides, quantum computation, diluted magnetic semiconductors, spintronic. Since 2001, he has been responsible of seven national super-computing projects (CINECA,CASPUR). Emiliano Bonera, Researcher Emiliano Bonera graduated at the Università degli Studi of Pavia (Italy) in 1998 after being also an exchange student at the University of Strathclyde in Glasgow. The topic of his dissertation was near-field microscopy. He obtained his PhD in 2002 from the University of Leeds with a thesis about 14

16 micro- and near-field optical characterisation of microelectronic materials. His interests span mainly on Raman and photoluminescence spectroscopy, internal photoemission spectroscopy, infrared spectroscopy, and near-field optical microscopy. He applied these techniques to the study of semiconductors, doped glasses, nanocrystals, high-κ materials, and other microelectronics-related issues. In the period he has been Post-Doctoral Fellow at the MDM Laboratory, where he has been working as a researcher from Sandro Ferrari, Researcher Sandro Ferrari is a research staff member at the MDM Laboratory. He received the Degree in Chemistry from Università degli Studi of Milano (Italy) in 1993 and the PhD from Università degli Studi of Brescia (Italy) in He joined the MDM Laboratory in 1998 and his reearch is focused on dielectrics for CMOS applications, organic and polymeric semiconductors for microelectronics and novel concepts in Non Volatile Memory devices. His main interest is the integration of polymers and oxides into micro- and nanoelectronic devices. Sandro Ferrari holds currently more than 50 publications in peerreviews journals, has organized two Workshops and co-ordinated the European Project ESQUI. He is currently the coordinator of the European IST project Versatile as well as of the PROTEO project founded by Fondazione Cariplo. Anna Maria Ferretti, Researcher Anna M. Ferretti is a researcher at the MDM Laboratory since July Her interests are focused on the characterization of semiconductors and oxide for microelectronic and quantum computing with CW and pulse EPR spectroscopy. She received the Degree in Chemistry in 1999 and the PhD in 2002 from Università degli Studi of Milano (Italy). During her PhD she spent one year at the Federal Institute of Technology of Zurich Switzerland at the ESR laboratory of Prof. Arthur Schweiger where she studied the pulse EPR technique. In 2002 she won the JEOL young investigator prize for her studies done on catalytic materials with pulse EPR. She spent her Post Doc (2003) at Fondazione San Raffaele with a Telethon grant studying proteins with NMR. Enrico Prati, Researcher Enrico Prati obtained his PhD for his work on microwave frequency transport in semiconductors at the Istituto per i Processi Chimico Fisici of the CNR and at the Università degli Studi of Pisa (Italy) in Part of the work was realized at Caswell Technology - Marconi Towcester, (UK). He joined the MDM Laboratory in The present research fields are both theoretical and experimental aspects of: electrically detected magnetic resonance, tunneling and localization in low dimensional electron systems under microwave irradiation, magnetic resonance effects of random telegraph noise in silicon MOS devices, and metamaterials. In February 2004 he has been awarded with the Young Scientist Award 2004 by the URSI Commettee for his work on negative refractive index propagation and metamaterials. He is now coordinating for MDM the project MARTA on metamaterials. He is presently Vice Secretary of ADI, the Italian PhD Society. Giovanna Scarel, Researcher Giovanna Scarel received the Laurea Degree in Physics from the Università degli studi of Trieste (Italy) and her PhD from the University of Wisconsin-Milwaukee (USA). From March 2003 she is reseracher at the MDM Laboratory. She is involved in the growth of high dielectric constant oxides relevant for advanced devices. Giovanna Scarel was the coordinator of a project on the ALD growth and characterization of rare earth oxides (PAIS-REOHΚ by INFM), and is presently working within the European Project REALISE on the same topic. Giovanna Scarel has more than 40 published papers and contributions to book chapters, and is editor of a book on rare earth oxides. 15

17 Sabina Spiga, Researcher Sabina Spiga graduated in Physics at the Università degli Studi of Bologna (Italy) in She obtained her PhD in Material Science from the Università degli Studi of Milano (Italy) in 2001 with a thesis on the synthesis and characterization of nanocrystals embedded in thin dielectric layers. In she was post-doc fellow at the MDM Laboratory; where she was appointed as research associate in Her current interest focuses on materials (high-κ oxides, oxides for resistive switching memories, nanocrystals) and processes for advanced CMOS devices, as well as on their electrical characterization. Her activity is carried out in the framework of European (GROWTH/NEON, IST/ET4US, IST/EMMA), Industrial and Italian R&D projects. Grazia Tallarida, Researcher Grazia Tallarida received the Laurea degree in Physics at the Università degli Studi of Roma (Italy) La Sapienza in After spending two years as reasearch assistant at the Engineering Department of the Cambridge University (UK), she joined MDM Laboratory in 96 where she currently holds a permanent position as researcher. At MDM Laboratory, she is in charge of the scanning probe microscopy activity and her main research topic is the development and application of SPM based techniques for the characterisation of thin films and nanostructures. Her interests include also the characterisation of the initial stages of growth of epitaxial oxides on semiconductors and the investigation of advanced materials for ultra-scaled devices. Claudia Wiemer, Researcher Claudia Wiemer graduated in Physics at the Università degli Studi of Milano (Italy) in She obtained her PhD in Applied Physics at the Ecole Polytéchnique Fédérale de Lausanne (Switzerland) in She joined Stanford University (USA) as a Post doctoral fellow in 1997, working on total reflection X-ray fluorescence analysis by synchrotron radiation. She joined MDM Laboratory in 2000, supported by a fellowship of the European project ESQUI. She has been a researcher associate at MDM Laboratory since Her main interests focus on the structural properties of oxide films, on the growth by MOCVD and characterization of calcogenides materials and on advanced X-ray methodologies. She is actively involved in the European projects ET4US, CHEMAPH (coordinator) and FOREMOST. She has currently 53 accepted refereed publications. Claudio Amabile, Post-Doctoral Fellow Claudio Amabile graduated in Physics at the Università degli Studi of Roma (Italy) La Sapienza in 2002 with a thesis about the dependence on doping of the resistivity tensor of high Tc superconductors in high magnetic fields determined through multicontact DC measurement. In his PhD thesis he studied the microwave response of various superconductors in high magnetic fields through a novel broadband technique; he also spent six months in the research center of Juelich (Germany), where he patterned and characterized magnesium diborides nanobridges for THz radiation receivers. He is currently working as postdoctoral researcher at MDM Laboratory where he is involved in the design and realization of metamaterials based microwave devices. His main interests are low temperature physics and electromagnetic behaviour of various systems. Dmitry Azamat, Post-Doctoral Fellow Dmitry Azamat obtained his PhD (Candidate of Sciences) in Solid State Physics from St. Petersburg State University in He joined the MDM Laboratory in 2005 to work on EPR of low-symmetry defects in bulk ZnO. His scientific interests include radiospectroscopy of point defects in solids, recombination processes in insulators and semiconductors, low-dimensional semiconductor structures, optical 16

18 detection of magnetic resonance. Currently he is research scientist at the A.F.Ioffe Physico-Technical Institute (Germany). Md. Nurul Kabir Bhuiyan, Post-Doctoral Fellow Md. Nurul Kabir Bhuiyan is a Post-Doctoral Fellow at MDM Laboratory since 2005, carrying out a research activity on the growth of high-κ oxide thin films on Ge, GaAs and Si substrates by MBE, and the characterization of films by RHEED, XPS and LEIS, in the framework of the European project ET4US. He obtained his Doctor of Engineering degree in Materials Science and Technology from Toyama University (Japan) in March 2005, where he mainly worked on the epitaxial growth of high-κ SrTi thin films on Si substrates by MBE and the characterization of films by RHEED, XPS, AFM and XRD. He received his B.Sc. (Honours) and M. Sc. in Physics, University of Dhaka (Bangladesh) in 1995 and 1997, respectively. Milena Georgieva, Post-Doctoral Fellow Milena Georgieva received the MSc degree in Solid State Physics in 1995 from the University of Sofia (Bulgaria). In January 2004 she obtained a PhD degree in Physics from the University of Salford (UK) for her work on thin magnetic films and trilayer devices with application in Spintronics. She has been a research scientist at the University of Sofia in the period and a post-doctoral fellow at the University of Salford between January August Since September 2005 she is a post-doctoral fellow at the MDM Laboratory. She gained experience in a number of techniques such as: VSM, SQUID magnetometry, AFM/MFM, XRD/XRR, TEM, photolithography and wet etching. Nolwenn Huby, post-doctoral Fellow Nolwenn Huby received the degree in Material Science in 2003 at the University of Rennes (France). In September 2006 she obtained her PhD in Electronics and Material Science from the University of Bordeaux (France). After the realisation of organic light emitting diodes (OLED), she performed electrical, optical and structural investigations in order to connect the molecular properties of the emissive molecules and the performances of the devices. Then she joined MDM Laboratory in October 2006 supported by a post-doctoral fellowship of the European project VERSATILE. The aim is the realization and the integration of non volatile memories based on II-VI and polymers semiconductors. Xiaolong Li, Post-Doctoral Fellow Xiaolong Li was a post-doctoral fellow at MDM Laboratory from September 2005 to September During this period, he mainly worked on the investigation of dielectric films, in the framework of the project SOLARIS concerning the characterisation of Magnetic Tunnel Junction. He obtained his PhD degree from the National Laboratory for Condensed Matter Physics, of the Chinese Academy of Sciences in The topic of his PhD thesis was the microstructural characterization of high-tc superconductors, high-κ dielectrics, and ferroelectric thin films. He has a rich experience and skills on X-ray scattering analysis, synchrotron radiation facility and XPS quantitative analysis. Hongliang Lu, Post-Doctoral Fellow Hongliang Lu received his PhD in Microelectronics from Fudan University of Shanghai (China) in June The topic of his dissertation was the investigation of high dielectric constant oxide films grown by atomic layer deposition (ALD). He won a fellowship at the MDM Laboratory in November His research interests are related to the growth and characterization of metal oxide thin films deposited by ALD. 17

19 Roberto Mantovan, Post-doctoral Fellow Roberto Mantovan graduated in 2002 at the Department of Physics at the Università degli Studi of Milano (Italy). In January 2006, he obtained his PhD from the same University presenting a thesis about the Mössbauer spectroscopy investigation of materials for non-volatile memory devices. R. Mantovan has been working at MDM Laboratory since His research activity is mainly focused on the application of Mössbauer spectroscopy to investigate the structural and magnetic properties of materials for spintronics. R. Mantovan is involved in on-line Mössbauer experiments at ISOLDE-CERN, in order to investigate the magnetic properties of Fe/Mn-implanted oxides and semiconductors. He works on the developing of lithographic process to fabricate structures and demonstrators for spintronics applications. Alessandro Molle, Post-Doctoral Fellow Alessandro Molle graduated in 2001 and obtained the PhD degree in Materials Science in 2005 at the Università degli Studi of Genova (Italy) after experiencing several beamtimes at the European Synchrotron Radiation Facility (ESRF) in Grenoble (France). He joined the MDM Laboratory in 2005 as a Post-Doc Fellow, involved in the molecular beam epitaxial growth and in situ characterization of high-κ oxide films on semiconductors, within the European Project ET4US. His Post-Doc fellowship has been recently renewed within the ATHOS Project in collaboration with IMEC (Belgium). Michele Perego, Post-Doctoral Fellow Michele Perego is post-doc fellow at the MDM Laboratory. He received the Degree in Physics from the Università degli Studi dell Insubria (Italy) in In February 2004 he obtained his PhD from the Università degli Studi of Milano (Italy) for his work on the TOF-SIMS characterization of nanocrystals embedded in thin oxide films. His research activity is focused on the synthesis and characterization of new materials (high-κ, nanocrystals) for advanced CMOS devices. He is expert in Secondary Ion Mass Spectrometry and X-ray photoelectron Spectroscopy. He holds more than 25 publications in peer reviews journals and he is currently the coordinator of the MATRIX project funded by Fondazione Cariplo. Inma Suàrez Lòpez, Post-Doctoral Fellow Inma Suàrez Lòpez graduated in Chemistry from the University of Málaga (Spain) in She completed her PhD at the University of Manchester in the UK under the supervision of Prof. J.C. Vickerman (Surface Analysis) and Prof. C.M. Carr (Textile Department). Her PhD thesis centered on the surface characterization and quantification of textile processes using ToF- SIMS. As an undergraduate she undertook several placements (Croatia - Pliva Co., University of Dortmund and Delft University) related to organic and analytical chemistry. Inma Suàrez Lòpez has been working at MDM Laboratory since Her current study focuses on novel routes for organic thin film transistors device fabrication by photolithography processes. Andrew Teren, Post-Doctoral Fellow Andrew Teren obtained his PhD in Materials Science and Engineering from Northwestern University, Chicago (USA) in Afterwards, he did post-doctoral work from at Forschungzentrum Juelich (Germany) on MOCVD growth and electrical characterization of high-κ dielectric thin films for microelectronic applications (within the European project Medea+). During , he worked at Belford Research Inc. (USA) on electrical characterization of strained Si devices. In he was a post-doctoral researcher at MDM Laboratory, where he was responsible for material characterization of chalcogenide materials and coordinator of the European project Chemaph of the 6 framework. 18

20 Stergios N. Volkos, Post-doctoral Fellow Stergios N. Volkos holds a BEng. (Hons.) degree in Electronics, MPhil. degree in Electrical Engineering and Electronics and PhD degree in Electrical Engineering and Electronics, all awarded by the University of Manchester (UK). Currently, he is a post-doctoral fellow in the MDM Laboratory working within the framework of an EU project called REALISE. His research field is predominantly focused on defects in high-κ dielectrics (transition metal and rare earth oxides) and at their interfaces with a silicon substrate. The experimental tools of primary interest include Current-Voltage (IV), Capacitance-Voltage (CV), Deep-Level-Transient-Spectroscopy (DLTS) and Inelastic- Electron-Tunneling-Spectroscopy (IETS) techniques. Giuseppe Bano, Post-Graduate Fellow Giuseppe Bano recived the Degree in Physics from the Università degli Studi of Milano (Italy) in December Topic of his dissertation was the investigation of point defect electronic structure in crystal oxide, by ab-initio Quantum Mechanical approach. In spring 2003, he attended classes in Nanotechnology at the Università degli Studi of Milano, and then won a fellowship in MDM Laboratory. In MDM his main activity concerned FT-IR characterization of thin high-κ oxide films and the developement of a smal ALD reactor for in-situ measurements. He is actually employed as Thin film Metrology Engineer in the R&D department of STMicroelectronics. Simone Malacrida, Post-Graduate Fellow Simone Malacrida received his Laurea Degree in Electronics Engineering from Politecnico di Milano (Italy) in The topic of his dissertation was the material composition of photosensitive optical fibers for Bragg gratings fabrication. He collaborated with Corning OTI (Italy) to develop a new fluorine core doped fiber (US Patent N ). After that, he worked at Underwriters Laboratories in industrial electronics safety certification. He joined MDM Laboratory in December His research activity was focused on the electrical properties of high-κ materials under ET4US European project. Concerning MAE bilateral project between Italy and Poland, he also collaborated with the Institute of Physics of the Polish Academy of Sciences (IFPAN) to perform Laplace DLTS on MIS devices for studying oxide interface defects. He has been also trained in experimental cryogenic processes. Francesca Perissinotti, Post-Graduate Fellow Francesca Perissinotti graduated in Material Science at the Università degli Studi of Torino (Italy) in 2003, with a thesis work about Morphological defects on Silicon Carbide for the production of electronic devices. In 2004 she started a post-graduate fellowship in MDM Laboratory to work on the process development for electronic devices, initially in the project Teseo Transistors with organic semiconductors and then in the project Proteo Electronic devices with n and p type organic semiconductors. During this period, she participated to the Organic Electronic Workshop in 2004 at IMEC (Leuven), to the MMD Meeting (Genova) and ICANS 21 (Lisbona) in Elisabetta Peron, Post-Graduate Fellow Elisabetta Peron graduated in Chemistry at the Università degli Studi of Padova (Italy) in 2001 and received her Master Course degree in Materials for Micro- and Nano-Technologies at the Università degli Studi of Pavia (Italy) in She worked at EniTecnologie, the corporate research center of Eni Gruop, where her main research activity was the characterization of nanoemulsions and of rocks saturated with complex fluid by NMR relaxometric technique. She joined MDM Laboratory in April 2004, where her main research activity was the study of polymeric semiconductors, regarding both their characterization and their use in devices. 19

21 Gabriele Seguini, Post-Graduate Fellow Gabriele Seguini is a post-graduate fellow at the MDM Laboratory. He received the Degree in Physics at the Department of Physics of the Università degli Studi of Milano (Italy) in His research activity is focused on the study of highκ material properties, in particular the band alignment, by means of Internal Photoemission Spectroscopy (IPE) and standard electrical characterization (CV, IV). He is also experienced in Inelastic Electron Tunneling Spectroscopy (IETS). He has been trained in some technological processes (thermal and e-beam evaporation) and cryogenics also in 3 He refrigerator. Silvia Baldovino, PhD Student Silvia Baldovino graduated in Material Science at the Università degli Studi of Torino (Italy) in In she worked at her PhD at the Walter Schottky Institut of München (Germany) concerning electron spin resonance investigations of silicon devices. She is currently working at the MDM Laboratory. She is mainly involved in the study of the defects present at the semiconductor/oxide interface and shallow donor in silicon, characterized with techniques based on the combination of standard electrical characterizations (IV, CV) with the electron spin resonance (EPR) spectroscopy. Luca Lamagna, PhD Student Luca Lamagna is a PhD student at the MDM Laboratory. He received his Bachelor of Science in Materials Engineering in October 2004 with a thesis about electrochemical deposition of thermoelectric nanostructured materials. In March 2006 he joined MDM Laboratory and he received his Master of Science in Materials Engineering in October 2006 with a thesis about ALD of Al 2 thin films on Si, Ge and GaAs. Since November 2006 he is a student of the PhD School in Nanostructures & Nanotechnologies at the Department of Materials Science of the Università degli Studi of Milano Bicocca (Italy). His research activity is focused on ALD growth and characterization (Spectroscopic Ellipsometry) of rare earth oxides, ALD reactors and processes; this activity is performed in the framework of the European Project REALISE. Jean-Luc Battaglia, Professor, Visiting scientist Jean-Luc Battaglia obtained his PhD in Mechanics and Engineering from the University of Bordeaux (France) in He joined the University of Bordeaux as a Lecturer in He directs the team of research Thermal Characterization of the TREFLE Laboratory (UMR CNRS 8508) at the Ecole Nationale Supérieures d Arts et Métiers since In the last years, he developed new approaches for the thermal characterization of thin films by modulated photothermal radiometry and time domain thermoreflectance. He obtained a chair of Professor at the University of Bordeaux in June His teaching activity mainly concerns heat transfer and fluid mechanics. He joined the MDM Laboratory as a visiting scientist in 2006 within the framework of the Intra-European Fellowship Marie Curie Project TCAMMD (August 2006-July 2007). During this project he was mainly interested by the thermal characterization of phase change material used for the non-volatile memory operating mode. He has currently 36 publications in peer-reviewed journals. Technical Staff Mario Alia, Technician Mario Alia is a process technician at the MDM Laboratory since February His main activities are: photolithography, process developing, installation and maintenance of the various equipments and upgrading the laboratory layout. His previous experiences were: - at Corning OTI: Process Engineer with projects responsibility to improve production profitability and costs, to manage issues in the layout and acquisition of clean room equipment ( ); - at Pirelli Optical Systems as responsible of the pilot line in the team of laser chip engineering ( ) and responsible of the Front End area: from the wafer with epitaxial structure to the chip realization ( ). He got the certification ISO 9002 for the manufacturing area under his responsibility; - at Pirelli Cable and Systems as responsible of a part of the IBM technological transfer of the 980 nm Pump Laser chip process ( ). 20

22 Simone Cocco, Technician Simone Cocco received the D.U. in Materials Science from the Università degli Studi of Cagliari (Italy) in In 2005 he attended courses on vacuum and leak detection issues at Varian Technology Inc. He is trained to work with various cryogenic systems (77K, 4K, 300mK). In 2007 he received the Laurea degree in Materials Science from the Università degli Studi of Cagliari. At MDM Laboratory he is in charge of Software and Hardware development for conventional and innovative measurements setups. He is involved in the realization of a growth chamber working in Atomic Layer Deposition regime allowing various non-conventional in-situ investigations. Roberto Colnaghi, Technician Roberto Colnaghi graduated in 2002 at the IT Department of Università degli Studi of Milano (Italy) discussing a project on software for mechanical engineering. He is now the system administrator and IT specialist of MDM Laboratory. He is involved in a project for the development of a web application for tracking the scientific and technical activity of the laboratory. Administrative Staff Anna Grazioli, Administrative staff Anna Grazioli graduated in Economics and Business at the Università degli Studi of Cagliari (Italy) in April She worked as administrative assistant at the Milan Research Unit from 2003 to 2005; since 2006 she s been working for MDM Laboratory. Mara Lanati, Scientific Secretary Mara Lanati, received the high school diploma in Technical Institute and Foreign Languages, Milano (Italy) in She is Scientific Assistant at the MDM Laboratory since June Her previous experiences are: : President Assistant and IT, R&D, Finance, Quality Assistant at Corning OTI SpA; : R&D Assistant at Pirelli Cavi e Sistemi SpA; : Raw Materials Qualification Area Assistant; : R&D Assistant at Pirelli SpA. Vania Zoccarato, Administrative Staff Vania Zoccarato holds a high School diploma in Accountancy and Foreign Languages, ITC M. Bassi, Seregno, (Italy) She worked as administrative assistant at the Milan Research Unit from 1999 to 2005; since 2006 she s been working at MDM Laboratory. Undergraduate Students Andrea Cappella, Università degli Studi of Trieste (Italy) Department of Physics, October May 2007; Omar Costa, Università degli Studi of Milano (Italy) Department of Physics, September November 2005; Roberto Fallica, Politecnico of Milano (Italy) Department of Electrical Engineering, December June 2007; Alessandro Freguglia, Università degli Studi of Milan Bicocca (Italy) Department of Materials Sciences, September April 2007; Johnny Curi Llamoja, Università degli Studi of Milano (Italy), Department of Physics, April June 2007; Cristiano Monguzzi, Università degli Studi of Milano (Italy), Department of Physics, July 2004 November 2005; Federica Orrù, Università degli Studi of Cagliari (Italy), Department of Materials Sciences, May November

23 Publications 2006 < A. Debernardi, M. Fanciulli, Structural and vibrational properties of high-dielectric oxides HfO 2 and TiO 2 : a comparative study, Mat. Sci. Semicon. Proc. 9, (2006). < M. Caymax, S. Van Elshotcht, M. Houssa, A. Delabie, T. Conard, M. Meuris, M.M. Heyns, A. Dimoulas, S.Spiga, M. Fanciulli, HfO 2 as gate dielectric on Ge: Interfaces and deposition techniques, Mat. Sci. Eng. B, 135, (2006). < R. Mantovan, C. Wiemer, A. Zenkevitch, M. Fanciulli, CEMS characterisation of Fe/high-κ oxide interfaces, Hyp. Int., 169, 1349 (2006). < M. Perego, G. Seguini, M. Fanciulli, Energy band alignment of HfO 2 on Ge, J. Appl. Phys., 100, (2006). < C. Dallera, F. Fracassi, L. Braicovich, G. Scarel, C. Wiemer, M. Fanciulli, G. Pavia, B. C. C. Cowie, Nondestructive diagnostics of high-κ dielectrics for advanced electronic devices, Appl. Phys. Lett, 89, (2006). < A. Molle, Md. N. K. Bhuiyan, G. Tallarida, M. Fanciulli, Formation and stability of germanium oxide induced by atomic oxygen exposure, Mat. Sci. in Semicond. Proc., 9, 673 (2006). < M. Fanciulli, E. Bonera, S. Nokhrin, G. Pacchioni, Phosphorous-oxygen hole centers in phosphosilicate glass films, Phys. Rev. B, 74, (2006). < G. Scarel, C. Wiemer, G. Tallarida, S. Spiga, G. Seguini, E. Bonera, M. Fanciulli, Y. Lebedinskii, A. Zekevich, G. Pavia, I.L. Fedushkin, G.K. Fukin, G.A. Domrachev, Atomic layer deposition of Lu silicate films using [(Me 3 Si) 2 N] 3 Lu, J. Electrochem. Soc., 153, F271 (2006). < S. Ferrari, S. Spiga, C. Wiemer, M. Fanciulli, A. Dimoulas, Germanium diffusion during HfO 2 growth on Ge by molecular beam epitaxy, Appl. Phys. Lett, 89, (2006). < C. M. Compagnoni, A. S. Spinelli, A. Bianchini, A. L. Lacaita, S. Spiga, G. Scarel, C. Wiemer, M. Fanciulli, Temperature dependence of transient and steady-state gate currents in HfO 2 capacitors, Appl. Phys. Lett., 89, (2006). < M. Perego, S. Ferrari, M. Fanciulli, Comparative study of negative cluster emission in sputtering of Si, Ge and their oxides, Appl. Surf. Sci., 252, 7236 (2006). < A. Molle, Md. N. K. Bhuiyan, G. Tallarida, M. Fanciulli, In situ chemical and structural investigation of the oxidation of Ge substrates by atomic oxygen, Appl. Phys. Lett., 89, (2006). < S. Ferrari, M. Fanciulli, Diffusion Reaction of Oxygen in HfO 2 /SiO 2 /Si Stacks, J. Phys. Chem. B., 110, (2006). < W. Porzio, S. Destri, M. Pasini, A. Bolognesi, A. Angiulli, P. Di Gianvincenzo, D. Natali, M. Sampietro, M. Caironi, L. Fumagalli, S.Ferrari, E. Peron, F. Perissinotti, 22

24 Oligo- and polymeric FET devices: Thiophene-based active materials and their interaction with different gate dielectrics, Mat. Sci. Eng. C, 26, 996 (2006). < E. Bonera, M. Fanciulli, G. P. Carnevale, Raman stress maps from finite-element models of silicon structures, J. Appl. Phys., 100, (2006). < A. Debernardi, Mn doped GaN/AlN Heterojunction for spintronic devices, Superlattices Microst., 40, 530 (2006). < E. Prati, M. Fanciulli, G. Ferrari, M. Sampietro, Effect of the triplet state on the random telegraph signal in Si n-mosfets, Phys. Rev. B, 74, (2006). < E. Prati, Microwave Propagation in Round Guiding Structures Based on Double Negative Metamaterials, Int. Journ. Infr. And Mill. Waves, 27, 1227 (2006). < S. Elliott, G. Scarel, C. Wiemer, M. Fanciulli, G. Pavia, Ozone-Based Atomic Layer Deposition of Alumina from TMA: Growth, Morphology, and Reaction Mechanism, Chem. Mat., 18, 3764 (2006). < A. Debernardi, A. Baldereschi, M. Fanciulli, Computation of the Stark Effect in P Impurity states in Silicon, Phys. Rev. B, 74, (2006). The paper was selected to appear in Virtual Journal of Quantum Information, 7 (2006), and in Virtual Journal of Nanoscale Science & Technology, 14, (2006). < M. Perego, M. Fanciulli, C. Bonafos, N. Cherkashin, Synthesis of mono and bi-layer of Si nanocrystals embedded in a dielectric matrix by e-beam evaporation of SiO/SiO 2 thin films, Mat. Sci. Eng. C, 26, 835 (2006). < A.V. Zenkevich, Y.Y. Lebedinskii, N.S. Barantsev, V.N. Nevolin, V.S. Kulikauskas, G. Scarel, M. Fanciulli, Degradation pattern of thin HfO 2 films on Si(100) under ultrahigh-vacuum annealing: an investigation by X-ray photoelectron spectroscopy, and low-energy ion scattering, Russian Microelectronics, 35, 210 (2006). < A. Debernardi, M. Fanciulli, First Principles Calculation of Materials for Spintronics and Nanoelectronic Devices, Science and supercomputing at CINECA, 2005-report, 453 (2006). < G. Seguini, S. Spiga, E. Bonera, M. Fanciulli, A. Reyes Huamantinco, C.J. Först, C.R. Ashman, P.E. Blöchl, A. Dimoulas, G. Mavrou, Band alignment at the La 2 Hf 2 O 7 /(001)Si interface, Appl. Phys. Lett., 88, (2006). < C. M. Compagnoni, A.S. Spinelli, A. Bianchini, A. L. Lacaita, S. Spiga,, M. Fanciulli, Characterization of transient currents in HfO 2 capacitors in the short timescale, Microelectronic Engineering, 83, 1927 (2006). < M. Perego, G. Seguini, G. Scarel, M. Fanciulli, X-ray photoelectron spectroscopy study of energy-band alignments of Lu 2 on Ge, Surf. Interface Anal., 38, 494 (2006). < C. M. Compagnoni, A. S. Spinelli, A. Bianchini, A. Lacaita, S. Spiga, M. Fanciulli, Transient currents in HfO 2 and their impact on circuit and memory applications, in Proc. IRPS, pp.124 (2006). 23

25 < M. Malvestuto, G. Scarel, C. Wiemer, M. Fanciulli, F. D Acapito, F. Boscherini, X-ray absorption spectroscopy study of Yb 2 and Lu 2 thin films deposited on Si(100) by atomic layer deposition, Nuclear Instruments and Methods in Physics, 246, 90 (2006). < P. Thibaudeau, A. Debernardi, Vinh Ta Phuoc, S. da Rocha, F. Gervais, Phonon anharmonicity in disordered MgAl 2 O 4 spinel, Phys.Rev.B, 73, (2006). < M. Alessandri, R. Piagge, S. Alberici, E. Bellandi, M. Caniatti, G. Ghidini, A. Modelli, C. Wiemer, S. Spiga, M. Fanciulli, High-κ materials in Flash Memories, ECS Transactions, 1, (2006). < G. Scarel, A. Svane, and M. Fanciulli, Scientific and technological issues of rare earth oxides: an introduction, in Rare earth oxide thin films: growth, characterization, and applications, Eds. M. Fanciulli and G. Scarel, Topics in Applied Physics vol Springer-Verlag (2006). < Yu. Lebedinskii, A. Zenkevich, G. Scarel, and M. Fanciulli, Film and interface layer composition of rare earth (Lu, Yb) oxides deposited by ALD, in Structural and compositional characterization of rare earth oxide films in Rare earth oxide thin films: growth, characterization, and applications, Eds. M. Fanciulli and G. Scarel, Topics in Applied Physics vol Springer-Verlag (2006). < S. Schamm, G. Scarel, and M. Fanciulli, Local structure, composition and electronic properties of rare earth oxide thin films studied using advanced transmission electron microscopy techniques (TEM-EELS), in Structural and compositional characterization of rare earth oxide films in Rare earth oxide thin films: growth, characterization, and applications, Eds. M. Fanciulli and G. Scarel, Topics in Applied Physics vol Springer-Verlag ( 2006). < S. Spiga, C. Wiemer, G. Scarel, O. Costa, M. Fanciulli, Electrical characterization of rare earth oxides grown by atomic layer deposition, in Rare earth oxide thin films: growth, characterization, and applications, Eds. M. Fanciulli and G. Scarel, Topics in Applied Physics vol Springer-Verlag (2006). < G. Seguini, M. Perego, M. Fanciulli, Experimental determination of the band offsets of rare earth oxides on different semiconductors, Rare earth oxide thin films: growth, characterization, and applications, Eds. M. Fanciulli and G. Scarel, Topics in Applied Physics vol Springer-Verlag (2006). < A.V Zenkevich, Yu. Lebedinskii, G. Scarel, and M. Fanciulli, XPS-LEIS study of high-κ rare earth (Lu, Yb) oxides and silicates on Si: the effect of annealing on microstructure evolution, in Defects in High-κ Gate Dielectric Stacks. NATO Science Series. II. Mathematics, Physics and Chemistry vol Ed. E. Gusev, Springer-Verlag (2006). < M. Fanciulli, O. Costa, S. Baldovino, S. Cocco, G. Seguini, E. Prati, and G. Scarel, Defects at the high-κ/semiconductor interfaces investigated by spin dependent spectroscopies, NATO Series book on Defects in High-κ Gate Dielectric Stacks, E. Gusev (ed.), Vol. 220, (2006). 24

26 2005 < M. Perego, S. Ferrari, M. Fanciulli, Negative cluster emission in sputtering of Si 1-x Ge x alloys: a full spectrum approach, Surf. Sci., 599, 141 (2005). < G. Ferrari, L. Fumagalli, M. Sampietro, E. Prati and M. Fanciulli, CMOS fully compatible microwave detector based on MOSFET operating in resistive regime, IEEE Microwave and Wireless Components Letters, 15, 7, 445 (2005). < C. Bonafos, N. Cherkashin, M. Carrada, H. Coffin, G. Ben Assayag, S. Schamm, P. Dimitrakis, P. Normand, M. Perego, M. Fanciulli, T. Muller, K. H. Heinig, A. Argawal, and A. Claverie, Manipulation of 2D arrays of Si nanocrystals by ultra-low-energy ion beam-synthesis for nonvolatile memories applications, Mater. Res. Soc. Symp. Proc. Vol. 830, (2005). < D. Skarlatos, C. Tsamis, M. Perego, M. Fanciulli, D.Tsoukalas, Interstitial injection during oxidation of very low energy nitrogen-implanted silicon, Mat. Sci. Eng B, 314, 124 (2005). < P. Fantini, A. Ghetti, C.P. Carnevale, E. Bonera, D. Rideau, A full self-consistent methodology for strain-induced effects characterization in silicon devices, IEDM Tech. Dig., 2005, 992 (2005). < A. Ferretti, M. Fanciulli, A. Ponti, A. Schweiger, Electron spin-echo relaxation and envelope modulation of shallow phosphorus donors in silicon, Phys. Rev. B, 72, (2005). < G. Brunoldi, S. Guerrieri, S. G. Alberici, E. Ravizza, G. Tallarida, C. Wiemer, T. Marangon, Self-annealing and aging effect characterization on copper seed thin films, Microelectronic Engineering, 82, 289 (2005). < R. Mantovan, S. Spiga, M. Fanciulli, Low temperature CEMS of Sn-implanted SiO 2, Hyp. Int., 165, 5 (2005). < E. Bonera, M. Fanciulli, Resonant Raman microscopy of stress in Silicon-based microelectronics, Microscopy of Semiconducting Materials XIV, 371, (2005). < S. Elliot, G. Scarel, C. Wiemer, M. Fanciulli, Y. Lebedinskii, A. Zenkevich, I.L. Fedushkin, Precursor combinations for ALD of rare earth oxides and silicates - A quantum chemical and X-ray study, Electrochem. Soc. Proc., 605, 2005 (2005). < E. Bonera, M. Fanciulli, M. Mariani, Raman spectroscopy of strain in subwavelength microelectronic devices, Appl. Phys. Lett., 87, (2005). < S. Spiga, C. Wiemer, G. Tallarida, G. Scarel, S. Ferrari, G. Seguini, M. Fanciulli, Effects of the oxygen precursor on the electrical and structural properties of HfO 2 films grown by atomic layer deposition on Ge, Appl. Phys. Lett., 87, (2005). < G. Ferrari, L. Fumagalli, M. Sampietro, E. Prati, M. Fanciulli, DC modulation in field-effect transistors operating under microwave irradiation for quantum readout, J. Appl. Phys., 98, (2005). < D. Skarlatos, C. Tsamis, M. Perego, M. Fanciulli, Oxidation-enhanced diffusion of boron in very low-energy N 2+ -implanted silicon, J. Appl. Phys., 97, (2005). 25

27 < K. Kukli, T. Aaltonen, J. Aarik, J. Lu, M. Ritala, S. Ferrari, A. Harsta, M. Leskela, Atomic Layer Deposition and Characterization of HfO 2 Films on Noble Metal Film Substrate, J. Electrochem. Soc, 152, f75 (2005). < S. Kremmer, H. Wurmbauer, C. Teichert, G. Tallarida, S. Spiga, C. Wiemer, M. Fanciulli, Nanoscale morphological and electrical homogeneity of HfO 2 and ZrO 2 thin films studied by conducting atomic-force microscopy, J. Appl. Phys., 97, (2005). < A. Chiodoni, C. Camerlingo, R. Gerbaldo, L. Gozzelino, F. Laviano, B. Minetti, C.F. Pirri, G. Rombolà, G. Tallarida, E. Tresso, E. Mezzetti, Transport characterization of silicon-ybco buffered multilayers deposited by magnetron sputtering, IEEE Trans. Appl. Supercond. 15, 3062 (2005). < D. Botta, C. Camerlingo, A. Chiodoni, F. Fabbri, R. Gerbaldo, G. Ghigo, L. Gozzelino, F. Laviano, B. Minetti, C.F. Pirri, G. Rombola, G. Tallarida, E. Tresso, E. Mezzetti, Intrinsic pinning and current percolation signatures in E-J characteristics of Si/YSZ/CeO 2 /YBCO layouts, Eur. Phys. J. B, (2005). < M. Malvestuto, R. Carboni, F. Boscherini, F. D Acapito, S. Spiga, M. Fanciulli, A. Dimoulas, G. Vellianitis, G. Mavrou, X-ray absorption study of the growth of Y 2 on Si(001), Phys. Rev. B, 71, (2005). < A. Debernardi, M. Peressi, A. Baldereschi, Spin polarization and band alignments at NiMnSb/GaAs interface, Comp. Mater. Sci., 33, 263 (2005). < E. Bonera, G. Scarel, M. Fanciulli, P. Delugas, V. Fiorentini, Dielectric Properties of High-κ Oxides: Theory and Experiment for Lu 2, Phys. Rev. Lett., 94, (2005). < E. Prati, M. Fanciulli, F. Capotondi, G. Biasiol, L. Sorba, A. Kovalev, J.D. Caldwell, C.R. Bowers, Magnetoresistively Detected Electron Spin Resonance in Low-Density Two-Dimensional Electron Gas in GaAs AlGaAs Single Quantum Wells, IEEE Trans. on Nanotechnology, 4, 100 (2005). < G. Scarel and M. Fanciulli, Thin-film engineering by atomic-layer deposition for ultra-scaled and novel devices, in Materials for Information Technology. Series on Engineering Materials and processes, Eds. E. Zschech, C.M. Whelan, and T. Mikolajick, Springer-Verlag London (2005) < A. Debernardi, Self-energy of phonons in semiconductors by density functional theory, Advances in Science and Technology, (2004). < G. Seguini, E. Bonera, S. Spiga, G. Scarel, M. Fanciulli, Energy band diagram of metal/lu 2 /silicon structures, Appl. Phys. Lett., 85, 5316 (2004). < M. Corti, S. Aldrovandi, M. Fanciulli, and F.Tabak, High-temperature correlated spin dynamics in FeSi from 29 Si NMR Relaxation, J. Magn. Magn. Mat., 112, 272 (2004). < E. Prati, G. Annino, M. Martinelli, Complex Variational Axial Mode Matching Method for Single-Mode and Overmoded Dielectric Resonators, Electromagnetics, 24, 8, 565 (2004). 26

28 < H. Schumann, I. L. Fedushkin, M. Hummert, G. Scarel, E. Bonera, M. Fanciulli, Crystal and Molecular Structure of [(η5-c 5 H 4 SiMe 3 ) 2 LuCl] 2 : A Precursor for the Production of Lu 2 Films, Z. Naturforsch. B, 59, 1035 (2004). < S. Ferrari, Characterization of High-κ Dielectrics with TOF-SIMS, Appl. Surf. Sci., , 609 (2004). < G. Weyer, H. P. Gunnlaugsson, K. Bharut-Ram, M. Dietrich, R. Mantovan, V. Naicker, D. Naidoo, R. Sielemann, Acceleration of diffusional jumps of interstitial Fe with increasing Ge concentration in Si 1-x Ge x alloys observed by Mössbauer spectroscopy, Hyp. Int., 158, 417 (2004). < M. Perego, S. Ferrari, M. Fanciulli, G. Ben Assayag, C. Bonafoss, M. Carrada, A. Claverie, Characterization of silicon nanocrystals embedded in thin oxide layers by TOF-SIMS, Appl. Surf. Sci., , 813 (2004). < G. Scarel, E. Bonera, C. Wiemer, G. Tallarida, S. Spiga, M. Fanciulli, I. L. Fedushkin, H. Schumann, Yu. Lebedinskii, A. Zenkevich, Atomic-layer deposition of Lu 2, Appl. Phys. Lett, 85, 630 (2004). < S. Ferrari, G. Scarel, Oxygen diffusion in Atomic Layer Deposited ZrO 2 and HfO 2 thin films, J. Appl. Phys., 96, 144 (2004). < K. Frohlich, K. Husekova, D. Machajdik, J. Hooker, M. Fanciulli, S. Ferrari, C. Wiemer, A. Dimoulas, G. Vellianitis, F. Roozeboom, Ru and RuO 2 electrodes for advanced CMOS technology, Mat. Sci. Eng. B, 109, 117 (2004). < S. Spiga, C. Wiemer, G. Tallarida, M. Fanciulli, M. Malvestuto, F. Boscherini, F. Dacapito, A. Dimoulas, G. Vellianitis, G. Mavrou, Structural characterization of epitaxial Y 2 on Si (001) and of the Y 2 /Si interface, Mat. Sci. Eng B, 109, 47 (2004). < G. Scarel, S. Spiga, C. Wiemer, G. Tallarida, S. Ferrari, M. Fanciulli, Trends of structural and electrical properties in atomic layer deposited HfO 2 films, Mat. Sci. Eng B, 109, 11 (2004). < S. Ferrari, M. Modrenau, G. Scarel, M. Fanciulli, X-ray reflectivity and spectroscopic ellipsometry as metrology tools for the characterization of interfacial layers in high-κ materials, Thin Solid Films, 450/1, 123 (2004). < E. Bellandi, A. Elbaz, S. Spadoni, R. Piagge, C. Coccorese, G. Pavia, S. Ferrari, S. Banerjee, M. Alessandri, Metrology issues in thin ONO stacks measurements by spectroscopic ellipsometry and X-ray reflectivity, Thin Solid Films, 450/1, 119 (2004). < C. Wiemer, S. Ferrari, M. Fanciulli, G. Pavia, L. Lutterotti, Combining grazing incidence X-ray diffraction and X-ray reflectivity for the evaluation of the structural evolution of HfO 2 thin films with annealing, Thin Solid Films, 450/1, 134 (2004). < L. Lutterotti, D. Chateigner, S. Ferrari, J. Ricote, Texture, residual stress and structural analysis of thin films using a combined X-ray analysis, Thin Solid Films, 450/1, 34 (2004). 27

29 < S. Banerjee, S. Ferrari, D. Chateigner, A. Gibaud, Recent advances in characterization of ultra-thin films using specular X-ray reflectivity technique, Thin Solid Films, 450, 23 (2004). < J. Ricote, D. Chateigner, M. Morales, M. L. Calzada, C. Wiemer, Application of the X-ray combined analysis to the study of lead titanate based ferroelectric thin films, Thin Solid Films, 450, 128 (2004). < K. Fršohlich, K. HušekovŽa, Z. Oszi, J.C. Hooker, M. Fanciulli, C. Wiemer, A. Dimoulas, G. Vellianitis, F. Roozeboom, Metal oxide gate electrodes for advanced CMOS technology, Ann. Phys-Berlin., 13, 31 (2004). < M. Perego, S. Ferrari, M. Fanciulli, G. Ben Assayag, C. Bonafos, M. Carrada, A. Clavarie, Detection and characterization of silicon nanocrystals embedded in thin oxide layers, J. Appl. Phys, 95(1), 257 (2004). < D. Skarlatos, E. Kapetanakis, C. Tsamis, M. Perego, S. Ferrari, M. Fanciulli, D. Tsoukalas, Oxidation of nitrogen-implanted silicon: comparison of nitrogen distribution and electrical properties of oxides formed by very low and medium energy N 2+ implantation, J. Appl. Phys.,96 (1), 300 (2004). < C. Bonafos, M. Carrada, N. Cherkashin, H. Coffin, D.Chassaing, G. Ben Assayag, A. Claverie, T. Muller, K.H. Heinig, M. Perego et al., Manipulation of two-dimensional arrays of Si nanocrystals embedded in thin SiO 2 layers by low energy ion implantation, J. Appl. Phys., 95(10), 5696 (2004). < A. Chiodoni, V. Ballarini, D. Botta, C. Camerlingo, F. Fabbri, S. Ferrari, R. Gerbaldo, G. Ghigo, L. Gozzelino, F. Laviano, B. Minetti, C. F. Pirri, G. Tallarida, E. Tresso, and E. Mezzetti, Characterization of Silicon-YBCO buffered multilayers grown by sputtering, Appl. Surf. Sci. 238, 485 (2004). < P. Normand, E. Kapetanakis, P. Dimitrakis, D. Skarlatos, K. Beltsios, D. Tsoukalas, C. Bonafos, G. Ben Assayag, N. Claverie, J.A. Van Den Berg at al, Nanocrystals manufacturing by ultra-low-energy ion-beam-synthesis for non-volatile memory applications, Nuclear Instruments & Methods in Physics Research, Section B (Beam Interactions with Materials and Atoms), 216, 228 (2004). < D. Skarlatos, C. Tsamis, D. Tsoukalas, M. Perego, S. Ferrari, M. Fanciulli, Nitrogen distribution during oxidation of low and medium energy nitrogen-implanted silicon, Nuclear Instruments & Methods in Physics Research, Section B (Beam Interactions with Materials and Atoms), 216, 75 (2004). < M. Perego, S. Ferrari, M. Fanciulli, G. Ben Assayag, C. Bonafos, A. Clavarie, Detection and characterization of silicon nanocrystals embedded in thin oxide layers, J. Appl. Phys., 95 (1), 257 (2004). < M. Fanciulli, S. Spiga, G. Scarel, G. Tallarida, C. Wiemer, G. Seguini, Structural and electrical properties of HfO 2 films grown by atomic layer deposition on Si, Ge, GaAs and GaN, Mat. Res. Soc. Symp. Vol 786, E6.14 (2004). 28

30 < E. Prati, Crossover between the cell size and wavelength of the incident radiation in a metamaterial, Microw. Opt. Techn. Lett., 40 no. (4), 269 (2004). < E. Prati, Microwave propagation in ferromagnetic semiconductors, J. Magn. Magn. Mater., vol.272, 3, 1999 (2004). < C. Wiemer, S. Ferrari, M. Fanciulli, G. Pavia, L. Lutterotti, Combining grazing incidence X-ray diffraction and X-ray reflectivity for the evaluation of the structural evolution of HfO 2 thin films with annealing, Thin Solid Films, 450, 134 (2004). Published Books Rare earth oxide thin films: growth, characterization, and applications, Eds. M. Fanciulli and G. Scarel, Topics in Applied Physics vol Springer-Verlag (2006). 29

31 Conferences 2006 < INVITED Talks E. Prati, Random Telegraph Signal in MOSFET devices, Hitachi Cambridge Laboratories; Cambridge (United Kingdom), November M.Fanciulli Shallow donors in silicon based nanostructures for quantum information processing: experiments and theory, 39 th Workshop: low-dimensional dynamical phenomena and simulations; Erice (Italy), July G. Scarel, S. Spiga, C. Wiemer, E. Bonera, G. Seguini, X.L. Li, and M. Fanciulli, Rare earth oxides by atomic layer deposition as gate dielectrics on Si and Ge: evaluation and perspectives Microelectronics meets Catalysis: innovative Oxide Materials Workshop; Hanse Wissenschaftkolleg, Delmenhorst (Germany), July M. Fanciulli High-κ dielectrics on high mobility substrates: interface properties, E-MRS 2006 Spring Meeting; Nice (France), May G. Scarel, S. Spiga, C. Wiemer, E. Bonera, G. Seguini, X.L. Li, M. Fanciulli, S. Schamm, I.L. Fedushkin, Yu. Lebedinskii, and A. Zenkevich, Rare earth oxides by atomic layer deposition as gate dielectrics on Si and Ge: evaluation and perspectives, E-MRS 2006 Spring Meeting; Nice (France), May M. Fanciulli, M.Perego, C. Bonafos, Nanocrystals in high-κ dielectric stacks for non-volatile memory applications CIMTEC; Acireale-Catania (Italy), June M. Fanciulli Spin coherence, manipulation, and detection in Si, Electron Spin Resonance and Related Phenomena in Low Dimensional Structures Workshop; Sanremo-Imperia (Italy), March < ORAL Presentations G. Scarel, C. Wiemer, S. Spiga, E. Bonera, G. Tallarida, S. Baldovino, M. Perego, G. Seguini, L. Lamagna, and M. Fanciulli, Atomic layer deposition: principle, applications, and precursors. Convegno su Nanomateriali: nuove opportunità per l Industria, CIA Chimica, Industria e Ambienti, Fiera di Brescia; Brescia (Italy), October M. Fanciulli, G. Scarel, S. Spiga, C. Wiemer, E. Bonera, G. Tallarida, X.L. Li, S. Baldovino, M. Perego, and G. Seguini, Atomic layer deposition of oxides for nanoelectronics and spintronics. XIII International Workshop on Oxide Electronics; Ischia Napoli (Italy), October G. Scarel, M. Perego, C. Wiemer, M. Fanciulli, I.L. Fedushkin, and A.A. Skatova, Fabrication of GeO 2 films by atomic layer deposition using a Ge(II)-based solid precursor. AVS 6 th International Conference on Atomic Layer Deposition; Seoul (Korea), July

32 S. Baldovino, S. Spiga, G. Scarel, and M. Fanciulli, Effect of the oxygen precursor on the interface between (100)Si and HfO 2 films grown by atomic layer deposition. 10th Europhysical Conference on Defects in Insulating Materials; Milano (Italy), July C. Wiemer, S. Spiga, E. Bonera, M. Fanciulli, R. Piagge, M. Alessandri, M. Caniatti, G. Pavia E. Cadelano, G. M. Lopez, V. Fiorentini, Structural and electrical properties of annealed Al 2 thin films as inter-poly dielectric for Flash Memory applications, WoDiM (workshop on dielectrics in microelectronics); Santa Tecla-Catania (Italy), June G. Scarel, X.L. Li, C. Wiemer, and M. Fanciulli, Atomic layer deposition of MnO and NiO using cyclopentadienyl (metallocene) - type of precursors, The Baltic conference on Atomic Layer Deposition, BALD2006; Oslo (Norway), June S. Baldovino, S. Spiga, G. Scarel, and M. Fanciulli, On the properties of the interface between (100)Si and HfO 2 films grown by ALD, E-MRS 2006 Spring Meeting; Nice (France), May E. Bonera, G. Scarel, S. Spiga, C. Wiemer, M. Fanciulli, Dielectric properties of La-based oxides deposited using atomic layer deposition, E-MRS 2006 Spring Meeting; Nice (France), May E. Bonera, M. Mariani, M. Fanciulli, Spectroscopy of strain in subwavelength microelectronics devices, E-MRS 2006 Spring Meeting; Nice (France), May A. Debernardi, M.Fanciulli, Structural and vibrational properties of high-dielectric oxides HfO 2 and TiO 2 : a comparative study, E-MRS 2006 Spring Meeting; Nice (France), May S. Ferrari and M. Fanciulli, Diffusion reaction of oxygen in HfO 2 /SiO 2 /Si Stacks, E-MRS 2006 Spring Meeting; Nice (France), May S. Ferrari, S. Spiga, C. Wiemer, M. Fanciulli, G. Mavrou and A. Dimoulas, Germanium diffusion during HfO 2 growth on Ge by MBE, E-MRS 2006 Spring Meeting; Nice (France), May A. Molle, M. N. K. Bhuiyan, G. Tallarida, and M. Fanciulli, In situ structural and chemical investigation on the oxidation of Ge(001) substrates by atomic oxygen, E-MRS 2006 Spring Meeting; Nice (France), May M. Perego, S. Spiga, M. Fanciulli, C. Bonafos, A. Mouti, S. Schamm, G. Benassayag, Nanocrystals in high-k dielectric stacks for non-volatile memory applications, E-MRS 2006 Spring Meeting; Nice (France), May G. Seguini, S. Spiga, E. Bonera, M. Fanciulli, S. Galata, A. Dimoulas, Band alignment at the MBE grown high-k oxides/semiconductor interface, E-MRS 2006 Spring Meeting; Nice (France), May S. Spiga, C. Wiemer, G. Scarel, M. Fanciulli, Electrical properties of rare earth oxide films grown by ALD, E-MRS 2006 Spring Meeting; Nice (France), May C.Wiemer, G. Scarel, M. Fanciulli, Crystallographic ordering induced by the substrate of HfO 2 films grown by ALD on Ge and GaAs, E-MRS 2006 Spring Meeting; Nice (France), May

33 E. Bonera, Ultra-scaled Non-Volatile Memories: The Problem of Mechanical Stress in the Techniological Nodes Beyond 60 nm, Expo 2006 Capitale Umano e Innovazione; Milano (Italy), March < POSTER Presentations A. Molle, M. N. K. Bhuiyan, G. Tallarida, and M. Fanciulli, In situ UHV experiments on Ge substrates: from surface preparation to thin film growth and characterization, ESF/PESC Exploratory Workshop Silicon/oxide Hetero-Epitaxy: A New Road/Towards a CMOS-Compatible Tunnel Diode Technology ; Como (Italy), September G.Tallarida, S.Spiga, C. Wiemer, M.Fanciulli, S.Kremmer, C.Teichert, Nanoscale electrical characterization of HfO 2 thin films by Conducting-AFM and Kelvin Probe Force Microscopy, ICN+T 2006; Basel, July A. Debernardi, Mn doped GaN/AlN Heterojunction for spintronic devices, E-MRS 2006 Spring Meeting (E-MRS - IUMRS - ICEM 06); Nice-France, May E. Prati, G. Ferrari, L. Fumagalli, M. Sampietro, M. Fanciulli, CMOS Compatible Microwave Power Detector Based on a Single MOSFET, III Simposio per le Tecnologie Avanzate Applicazione delle Nanotecnologie per la Difesa nei settori Strutturale, Elettronico, Biotecnologico ; Roma (Italy), June E. Prati, M. Fanciulli, F. Capotondi, G. Biasiol, L. Sorba, A. Kovalev, J. D. Caldwell, and C.R. Bowers, Magnetoresistively Detected ESR of Defects in a GaAs/AlGaAs Heterostructure, Electron Spin Resonance and Related Phenomena in Low Dimensional Structures Workshop; Sanremo-Imperia (Italy), March A. Debernardi, A. Baldereschi, and M.Fanciulli, Theory of Shallow Impurities in Si for Quantum Computing, Electron Spin Resonance and Related Phenomena in Low Dimensional Structures Workshop; Sanremo-Imperia (Italy), March E. Prati, M. Fanciulli, G. Ferrari, M. Sampietro, Random Telegraph Noise in Si MOSFETs: Towards Single Spin Resonance Read-out in Quantum Devices, Mauterndorf Winter School; Mauterndorf (Austria), February < INVITED Talks C. Wiemer, X-ray characterization at the MDM Laboratory INFM workshop: X-ray reflectivity measurements for the evaluation of thin films and multilayers thickness, First Workshop of the VAMAS project; Brescia (Italy), September M. Fanciulli, Defects at the high-κ/semiconductor interfaces investigated by spin-dependent spectroscopies, NATO Advanced Research Workshop Defectsin Advanced High-κ Dielectrics Nano-Electronic Semiconductor Devices; St. Petersburg (Russia), July M. Fanciulli, More Moore: opportunities for fundamental and applied research, EARMA Conference; Genova (Italy), June M. Fanciulli, High-κ dielectrics for ultra-scaled and emerging nanoelectronic devices, MMD Meeting 2005; Genova (Italy), June

34 M.Fanciulli, E. Prati, G. Ferrari and M. Sampietro, Random Telegraph Signal in n-mosfets: a way toward Single Spin Resonance Detection ; UPON Gallipoli-Lecce (Italy), June M. Fanciulli, Electrical characterization of rare earth oxides and of their interfaces with semiconductors, ESF Exploratory Workshop on: Rare earth oxide thin films: growth, characterization, and applications ; Sanremo-Imperia (Italy), May G. Scarel and M. Fanciulli, Scientific and technological issues of rare earth oxides: introduction to the main topics of interest of the Workshop. ESF Exploratory Workshop: Rare earth oxide thin films: growth, characterization, and applications; Sanremo-Imperia (Italy), May G. Seguini, Experimental determination of the band offsets of rare earth oxides on different semiconductors, ESF Exploratory Workshop on: Rare earth oxide thin films: growth, characterization, and applications ; Sanremo-Imperia (Italy), May M. Fanciulli Atomic layer deposition of high-κ dielectrics on Ge and GaAs 2005 MRS Spring Meeting; San Francisco (California-USA), March E. Bonera, Raman Spectroscopy for Strain Characterisation of Microelectronic Devices, Seminars at ETH Zurich (Switzerland), March M. Fanciulli, Opportunities for fundamental and applied research, Casimir Workshop; University of Milano, Milano (Italy), February < ORAL Presentations E. Prati, G. Ferrari, M. Sampietro, P. Fantini, M. Fanciulli, Microwave Induced Effects on the Random Telegraph Signal in a MOSFET, UPON; Gallipoli-Lecce (Italy), June E. Prati, M. Fanciulli, E. Dalle Mese, Metamateriali con Passo Reticolare Confrontabile con la Lunghezza d Onda, II Simposio per le Tecnologie Avanzate Applicazione delle Nanotecnologie per la Difesa nei settori Strutturale, Elettronico, Biotecnologico ; Roma (Italy), June E. Prati, M. Fanciulli, G. Ferrari, M. Sampietro, High Magnetic Field Dependence of Capture/Emission Fluctuations of a Single Defect in Silicon MOSFETs, ICNF 2005; Salamanca (Spain), September M. Perego, G. Seguini, G. Scarel, M. Fanciulli, X-ray photoelectron spectroscopy measurements of band alignment of rare earth oxides ECASIA 05; Vienna (Austria), September G. Scarel, C. Wiemer, S. Spiga, E. Bonera, M. Fanciulli, I. L. Fedushkin, Yu. Lebedinskii, and A. Zenkevich, Effects of precursor and substrate choice on the properties of thin Lu 2 and Yb 2 films deposited using atomic layer deposition. AVS 5 th International Conference on Atomic Layer Deposition; San José (California-USA), August G.Tallarida, S. Spiga, C. Wiemer, M. Fanciulli, S. Kremmer, H. Wurmbauer, C. Teichert, Nanoscale electrical properties of high-κ dielectrics by scanning probe methods, MMD meeting; Genova (Italy), June

35 E. Bonera, M. Mariani, M. Fanciulli, Raman Mapping of Stress in Subwavelength Microelectronic Devices MMD meeting; Genova (Italy), June S. Spiga, G. Scarel, C. Wiemer, G. Tallarida, S. Ferrari and M. Fanciulli, Effects of the oxygen precursors on the electrical and structural properties of HfO 2 films grown by ALD on Ge, MRS Spring Meeting; San Francisco (California-USA), March E. Bonera, M. Fanciulli, Resonant Raman Microscopy of Stress in Silicon Based Microelectronics, Microscopy of Semiconducting Materials XIV; Oxford (United Kingdom), April C. Wiemer, G. Tallarida, S. Ferrari, M. Fanciulli, S. Kremmer, C. Teichert, J.W. Seo, C. Dieker, A. Dimoulas, Characterization of MBE growh HfO 2 films on Ge(001), MRS Spring Meeting; San Francisco (California-USA), March E. Bonera, G. Scarel, M. Fanciulli, P. Delugas, V. Fiorentini Dielectric Properties of High-κ Oxides Theory and Experiment for Lu 2 American Physical Society; Los Angeles (California- USA), March < POSTER Presentations R. Mantovan, C. Wiemer, Z. Zenkevich, and M. Fanciulli, CEMS characterization of Fe/high-κ oxide interfaces, International conference on the application of the Mössbauer effect ICAME 2005; Montpellier (France), September R. Mantovan, S. Spiga, and M. Fanciulli, Low temperature CEMS of Sn-implanted SiO 2 International conference on the application of the Mössbauer effect ICAME 2005; Montpellier (France), September M. Perego, S. Ferrari, M. Fanciulli, Comparative study of negative cluster emission in sputtering of Si, Ge, and their oxides, SIMS XV 15 th International Conference on Secondary Ion Mass Spectrometry; Manchester (United Kingdom), September E. Prati, M. Fanciulli, G. Ferrari, M. Sampietro, P. Fantini, Microwave Induced Effects on the Random Telegraph Signal in a MOSFET, ICNF 2005; Salamanca (Spain), September G. Scarel, S. Spiga, E. Bonera, C. Wiemer, G. Tallarida, G. Seguini, M. Fanciulli, I.L. Fedushkin, Yu. Lebedinskii, and A. Zenkevich, Rare earth oxides for microelectronic, nanoelectronic, and spintronic applications, MMD Meeting; Genova (Italy), June S. Ferrari, E. Peron, F. Perissinotti, G. Scarel, D. Natali, M. Caironi, L. Fumagalli, M. Sampietro, A. Angiulli, P. Di Gianvincenzo, A. Bolognesi, The role of Al 2 and SiO 2 gate dielectric surfaces on charge transport phenomena in P3HT based transitori, ICANS; Lisbon (Portugal), June R. Mantovan, C. Wiemer, Z. Zenkevich, and M. Fanciulli, CEMS characterization of Fe/high-κ oxide interfaces, Nanomagnetism and Spintronics spring school; Cargese (France), June S. Spiga, G. Scarel, C. Wiemer, G. Tallarida, G. Seguini, S. Ferrari, M. Perego, and M. Fanciulli, Atomic layer deposition of high-dielectric constant oxides on Ge and GaAs for ultra-scaled devices, MMD Meeting; Genova (Italy), June

36 C. Wiemer, S. Spiga, G. Tallarida, S. Ferrari, G. Seguini, A. Molle, M. Fanciulli M. Seo, C. Dieker, A. Dimoulas, High-κ oxides grown by molecular epitaxy on high mobility substrates, MMD Meeting; Genova (Italy), June F. Wallrapp, P. Fromherz, G. Scarel, S. Spiga, C. Wiemer, G. Tallarida, and M. Fanciulli, High-κ oxides deposited by ALD for neuroelectronic applications, MMD Meeting; Genova (Italy), June M. Perego, S. Spiga, M. Fanciulli, C. Bonafos, N. Cherkashin, Synthesis of mono and bi-layer of Si nanocrystals embedded in a dielectric matrix by e-beam evaporation of SiO/SiO 2 thin films, E-MRS spring meeting; Strasbourg (France), May S. Ferrari, E. Peron, F. Perissinotti, G. Scarel, D. Natali, M.Caironi,L. Fumagalli, M. Sampietro, A. Angiulli, P. Di Gianvincenzo, A. Bolognesi, Organic Thin Film Transistors obtained using Atomic Layer Deposited Al 2 as high-κ dielectric, MRS Spring meeting; S.Francisco (California-USA), March < INVITED Talks G. Scarel and M. Fanciulli, Atomic layer deposition of rare earth oxide thin films for ultra-scaled and novel devices, XC Congresso Nazionale della Società Italiana di Fisica; Brescia (Italy), September M. Fanciulli, Shallow donor elecron spins as qubits in silicon: decoherence and hyperfine interaction manipulation, 2004 IEEE NTC Quantum Device Technology Workshop, Clarkson University; New York (USA) May E. Bonera, Optical Characterisation of Nano-Aggregates, seminars at Università di Milano Bicocca; Milano (Italy), May E. Prati, Metamaterials Engineering the Emerging Properties of Materials (From Negative Refractive Index to Microwave Technology), Università di Napoli Federico II; Napoli (Italy), February < ORAL Presentations M. Perego, S. Ferrari, M. Fanciulli, Negative cluster emission in sputtering alloys: a full spectrum approach, SIMS Europe 2004; Muenster (Germany), September G. Scarel, S. Spiga, C. Wiemer, G. Tallarida, E. Bonera, G. Bano, G. Seguini, S. Baldovino, M. Fanciulli, I.L. Fedushkin, H. Schumann, Y. Lebedinskii, and A. Zenkevich, Atomic layer deposition of Lu 2 using H 2 O and as oxygen precursors, ALD2004 Congress; Helsinki (Finland), August S. Ferrari, G. Scarel and M. Fanciulli Mass transport phenomena n HfO 2, 16 th International Vacuum Congress; Venezia (Italy), June G. Scarel, S. Spiga, G. Seguini, C. Wiemer, G. Tallarida, S. Baldovino, E. Bonera, M. Fanciulli, I.L. Fedushkin, H. Schumann, Y. Lebedinskii, and A. Zenkevich, Atomic Layer Deposition grown rare earth oxides for gate dielectric applications, 16 th International Vacuum Congress; Venezia (Italy), June

37 E. Prati, J. Ravazzola, Merging and Crossing Modes in Double Negative Metamaterials, URSI 2004; Pisa (Italy), May G. Scarel, S. Spiga, C. Wiemer, G. Tallarida, S. Baldovino, E. Bonera, G. Bano, S. Ferrari, and M. Fanciulli, Influence of oxygen precursors (H 2 O or ) on the properties of as grown and annealed HfO 2 films grown by atomic layer deposition from HfCl 4 and Hf(O t Bu) 2 (mmp) 2, E-MRS 2004 Spring Meeting; Strasbourg (France), May G. Scarel, S. Spiga, G. Seguini, C. Wiemer, G. Tallarida, E. Bonera, M. Fanciulli, I.L. Fedushkin, H. Schumann, A. Zenkevich, and Y. Lebedinskii, ALD grown lutetium-based oxides for gate dielectric applications, 2004 MRS Spring Meeting; San Francisco (California-USA), April M. Perego, S. Ferrari, M. Fanciulli, A. Claverie, C. Bonafos, M. Carrada, G. Ben Assayag, D. Chassaing, P. Normand, P. Dimitrakis, E. Kapetenakis, T. Muller, K. H. Heinig, V. Soncini, 2D-arrays of Si nanocrystals embedded in thin SiO 2 layers for new memory devices, V Silicon Workshop 2004; Genova (Italy), February G. Scarel, E. Bonera, C. Wiemer, G. Tallarida, S. Spiga, S. Ferrari, M. Fanciulli, I.L. Fedushkin, H. Schumann, Y. Lebedinskii, and A. Zenkevich, Lu 2 films by atomic layer deposition: growth and characterization, V Silicon Workshop 2004; Genova (Italy), February C.Wiemer, S. Spiga, S. Ferrari, G. Seguini, G. Tallarida, M. Fanciulli, A. Dimoulas, G. Apostolopoulos, G. Vellianitis, J. C. Hooker, Z. M. Rittesma, Amorphous and Cristalline La 2 Hf 2 O 7 for gate dielectric applications, V Silicon Workshop 2004; Genova (Italy), February A. Debernardi, M. Peressi, A. Baldereschi, Maintaining half-metallicity at NiMnSb/III-V semiconductors: the role of semiconductor substrate and local interface termination, XC Congresso nazionale società italiana di fisica; Brescia, (Italy), September A. Debernardi, A. Baldereschi, M. Fanciulli, Electric-field dependence of the ground-state hyperfine splitting of P impurities in silicon, XC Congresso nazionale società italiana di fisica; Brescia, (Italy), September < POSTER Presentations E. Bonera, G. Scarel, M. Fanciulli, Infrared Characterisation of High-κ Materials, INFMeeting; Genova (Italy), June G. Scarel, C. Dallera, L. Braicovich, F. Fracassi, M. Fanciulli, C. Wiemer, and B.C. Cowie, In depth characterization of high-κ dielectrics for ultra scaled CMOS devices by hard X-ray photoemission, INFMeeting; Genova (Italy), June G. Scarel, S. Spiga, G. Seguini, C. Wiemer, G. Tallarida, S. Baldovino, E. Bonera, M. Fanciulli, I.L. Fedushkin, H. Schumann, Y. Lebedinskii, and A. Zenkevich, ALD grown rare earth oxides for gate dielectric applications, INFMeeting; Genova (Italy), June C. Wiemer, S. Ferrari, C. Marchiori, S. Spiga, G. Seguini, G. Tallarida, E. Bonera, M. Fanciulli, G. Norga, J. Fompeyrine, J.-P. Locquet, G. Apostolopoulos, A. Dimoulas, 36

38 Epitaxial and amorphous high-κ oxides grown by molecular beam epitaxy, INFMeeting; Genova (Italy), June C.Wiemer, C. Marchiori, G. Scarel, S. Spiga, S. Baldovino, S. Ferrari, G. Pavia, M. Fanciulli, Properties of the interfacial layer between silicon and HfO 2 films deposited by atomic layer deposition using different precursor combinations, (Workshop on Dielectrics in Microelectronics); Kinsale-Cork (Ireland), June E. Prati, M. Fanciulli, F. Capotondi, G. Biasiol, L. Sorba, A. Kovalev, J. D. Caldwell, and C.R. Bowers, Magnetoresistively Detected Electron Spin Resonance in Low Density Two Dimensional Electron Gas in GaAs/AlGaAs Single Quantum Wells, Quantum Device Technology Workshop; Potsdam (New York - USA), May Organization of conferences ø International Workshop on Electron spin resonance and related phenomena in low dimensional structures 6-8 March 2006, San Remo - Imperia (Italy) Convened by: Marco Fanciulli and Enrico Prati ø ESF PESC Exploratory Workshops Rare Earth Oxide Thin Films: growth, characterization, and applications May 2005, San Remo - Imperia (Italy) Convened by: Marco Fanciulli and Giovanna Scarel 37

39 Research Projects INFM Projects < PA REOHK - Rare Earth Oxides as High-κ ( ). Objectives. The main objectives of this project are (1) the growth by ALD of Yb 2 and Lu 2 films, (2) their structural, physical, chemical, electrical, electronic, and interface characterization and (3) the understanding of their physical properties using a number of techniques including also specialized methods such as X-ray absorption techniques and Density Functional Theory calculations. Yb 2 and Lu 2 films have a potential for applications as gate and interpoly dielectrics in ultra scaled CMOS devices. Candidates to be implemented as gate dielectrics in CMOS devices should possess a high dielectric constant (κ), large valence and conduction band offsets (VBO, CBO). Large band gap, high thermodynamical stability against silicate and silicide formation, low defects and interfacial states densities are also required. For the Rare Earth Oxides, the theoretical as well as experimental information on these physical quantities are scarce or absent. Nevertheless, the limited knowledge on these materials suggests that rare earth oxides are indeed promising candidates: we aim at demonstrating this experimentally and theoretically. The second goal of the project is thus to assess theoretically and measure experimentally the physical quantities bearing on their applicability as gate oxides in CMOS devices. Partners. INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy; UdR Cagliari, Italy; UdR Bologna, Italy. Role. Coordinator (G. Scarel) Industrial research projects <MDM-STM Advanced materials and characterization for nm technology nodes non-volatile memory devices. Objectives. The planned research activity is related to the growth and characterization of materials for advanced microelectronic devices, mainly non-volatile memories, of interest for STMicroelectronics (STM) and it is based on the requirements for the technology nodes beyond the 65 nm with reference to the Front End Processes, Interconnect, and Emerging Research Devices sections of the International Semiconductor Roadmap for Semiconductors, 2003 Edition. In particular the following topics are covered: High-κ dielectrics for NVM; Mapping of stress in silicon, Low-κ materials characterization: SiOC and BPSG-PSG; Characterization of materials for interconnections, Characterization of stress induced defects in tunnel oxide (SiO 2 based) for NVM, Materials for Phase Change Memories; Oxidation kinetics of Silicon in non-standard conditions. Coordinator: M. Fanciulli <MDM-STM Advanced materials and characterization for the nm technology node non-volatile memory devices. Objectives. The planned research activity is related to the growth and characterization of materials for advanced microelectronic devices, mainly non-volatile memories, of interest for STMicroelectronics (STM). The activity is coherently linked with all the other activities and expertise of the MDM-INFM National Laboratory. The planned activity takes into consideration the specific needs of STM and is based on the requirements to reach the 45 nm technology node with reference to the Front End Processes, Interconnect, and Emerging Research Devices sections of the International Semiconductor Roadmap for Semiconductors, 2003 Edition. In particular the following topics are covered: Evaluation of Hf-based ternary oxides deposited at MDM using different precursors; Characterization 38

40 of high-κ dielectrics for NVM; Development of a process for the realization of test devices; Mapping of stress in silicon; Low-κ materials characterization: SiOC; Characterization of materials for interconnections; Materials for Phase Change Memories; Characterization of stress induced defects in tunnel oxide (SiO 2 based) for NVM. Coordinator: M. Fanciulli <MDM-STM Advanced materials and characterization for the 45 nm technology node non-volatile memory devices and beyond. Objectives. The planned activity takes into consideration the specific needs of STM and is based on the requirements to reach the 45 nm technology node with reference to the Front End Processes, Interconnect, and Emerging Research Devices sections of the International Semiconductor Roadmap for Semiconductors, 2003 Edition. In particular the following topics are covered: ALD and characterization of high-κ dielectrics for NVM; Characterization of high-κ dielectrics for NVM deposited by STM; Mapping of stress in Si; Low-κ materials characterization; Characterization of materials for interconnects; Materials for Phase Change Memories; Oxides for resistive switching NVM. Coordinator: M. Fanciulli National research projects <MIUR FIRB - Quantum phases of ultra-low electron density semiconductor heterostructures ( ). Objectives. Investigation of ultralow density two-dimensional electron gasses (2DEG) confined in high-mobility AlGaAs/GaAs heterostructures and few-electrons state confined in AlGaAs/GaAs artificial quantum dots. Advancements in this area require the application and optimization of the most advanced semiconductor growth protocols and nanofabrication techniques and the use of a large spectrum of experimental and theoretical techniques for the study of system s properties. These activities are at the core of this project. Attention will be given in particular to the spin properties of these states, to signatures of possible crystalline order and to the identification of quantum phase transitions between ground states with different magnetic properties. To this end specific efforts are directed to the optimization of experimental protocols for the observation of spin and charge properties of such systems down to the ultimate limit of single-electron detection. The project builds on an international collaboration between experimental groups at TASC, Trieste (MBE growth of high-mobility AlGaAs/GaAs heterostructures), Scuola Normale Superiore SNS, Pisa (nanofabrication, magneto-optics and magneto-transport), MDM, Milano (electron spin resonance studies), Bell laboratories/columbia University, USA (MBE growth and optical spectroscopy), and the theoretical group in Modena (theory of few-electron confined systems). Partners: University of Modena (coordinator); Scuola Normale Superiore Pisa; Laboratorio TASC-INFM; INFM, Laboratorio Nazionale MDM. Role: Partner. MDM responsible: M. Fanciulli <MIUR FIRB - Semiconductor/Superconductor structures for an integrated electronics ( ). Objectives. The present project is based on the very promising results obtained in the INFM-PAIS DEBUSSY (Sez.E-Sez.D), aimed at the deposition and characterization of YBa 2 Cu 3 O 7-x (YBCO) films on buffered Si. This project represents the next stage of the research, aiming to reach a real control, reproducibility and optimization of the multilayered YBCO/buffer/Si structures, to achieve a complete understanding of the fundamental physics involved, and to obtain superstructures suitable for effective miniaturized micro-integrated semiconductor-superconductor devices. Moreover, innovative solutions for the cooling and temperature control by means of integrated thermal bus are proposed. 39

41 Partners. Politecnico Torino (coordinator); CNR-IC; ENEA; University of Roma 3; University of Torino; INFM Laboratorio Nazionale MDM. Role. Partner. MDM responsible: G. Tallarida. <MIUR FIRB - Miniaturized systems for electronics and photonics ( )- FIRB RBNE012N3X. WP2: Methods for the scaling of CMOS devices Activity 4: Advanced dielectrics for ultra-scaled CMOS devices Objectives. The MDM contribution is related to the growth and characterization of advanced high-κ dielectrics (HfO 2 ) grown using ALD for CMOS technology nodes below 65 nm. The materials and structures produced by the MDM Laboratory will be also supplied to some of the partners (Politecnico di Milano, Università di Roma 1 La Sapienza,) for specific additional characterizations. Partners. INFM Laboratorio Nazionale MDM; Politecnico di Milano; Università di Roma1 La Sapienza. Role. Partner. MDM responsible: M. Fanciulli. <CINECA - Spintronic properties of Mn-based semiconductor heterostructures ( ). Objectives. Within the project we have performed a systematic study of different mechanisms acting on polarized electrons at the interfaces Heusler compound/iii- V semiconductor. Our target is to give an important contribution to the individuation of junctions suitable to realize spintronic devices. In particular, they must allow the injection of polarized spin in semiconductors. For this reason we have computed the band allignement of epitaxial NiMnSb on three different zinc-blende substrates (001)- oriented: GaAs, InP, and GaSb, in order to undertand how the lattice parameter and the chemical composition of the subsrate influence the majority spin Schottky barriers and the minority spin band offset of these interfaces. Partners. INFM Laboratorio Nazionale MDM; CRS Democritos (Trieste); Dip. Fisica Teorica Università di Trieste; EPFL-Lausanne, Switzerland. Role. Coordinator (A. Debernardi). <CINECA - First principles calculation of materials for spintronics and nanoelectronic devices (2006). Objectives. The objective is the study of the band alignment in GaN/AlN:Mn heterostructures to identify a junction capable to inject spin-polarized electrons in a semiconductor. We predicted that this heterostructure exibit half-metallic properties, that make it suitable as a spin injector. Further, we have studied the doping of semiconducting (or insulating) oxides with magnetic elements to find materials that are both semiconductor and ferromagnetic at room temperature and to understand how their magnetization is effected by the presence of defects. The final target is the design of a spintronic device whose properties can be tailored according to the needs of electronic industry. The result are compared experimental data obtained at CERN within the ISOLDE collaboration. Partners. CNR-INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy; ISOLDE- CERN, Geneve, Switzerland. Role. Coordinator (A.Debernerdi) <CASPUR (CINECA) - Electronic and magnetic properties of Mn-based materials and nanostructures (2006). Objectives. We have proposed an extensive study of electronic, structural, and magnetic properties of different Mn-based heterostructures. We have computed by planewave pseudo-potentials method the magnetization of Mn nano-structure on different ferromagnetic substrate to integrate and complete already performed tight binding calculations. In particular, we have focused our effort on Mn and Mn-alloy epilayers grown Ni substrates with different orientationis, namely Mn/Ni(001), Mn/Ni(110), Mn/ 40

42 Ni(111) that have recently attracted considerable interest by experimentalists for their unusual magnetic properties. Our goal is the understand of microscopic mechanisms responsible of alloying and magnetism of these Mn-nanostructures. Partners. CNR-INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy; International Center of Theoretical Physics (ICTP) Trieste, Italy; CNRS-Strasburg, France; Univ. Brazzaville, Congo. Role. Coordinator (A. Debernardi). <CARIPLO TESEO - Sviluppo di Tecnologie a Semiconduttori Organici per Applicazioni Optoelettroniche ( ). Objectives. Although known since a long time, it is only recently that polymeric semiconductors are becoming of extreme interest for a wide variety of application such as Organic Thin Film Transistors, Organic Light Emitting Diodes and Organic Photovoltaic Cells. This project aims to explore novel functional polymer and novel routes of device fabrication that couple standard microlectroncic techniques such as optical photolithography with the use of organic semiconductors. Partners. CNR-INFM, Laboratorio Nazionale MDM, Agrate Brianza; CNR-ISMAC; Politecnico di Milano. Role. Coordinator (S. Ferrari). <CARIPLO PROTEO - PolimeRi e molecole Organiche per Tecnologie Elettroniche ed Optoelettroniche ( ). Objectives. This project was born as a continuation of the positive experience carried out during the TESEO project and aims at the fabrication of n-type organic transitors, and OLED with Infrared Emission, and explores the application of photodiodes, as detectors in a wide range of light spectrum. Parteners. CNR-INFM, Laboratorio Nazionale MDM, Agrate Brianza; CNR-ISMAC; Politecnico di Milano. Role. Coordinator (S. Ferrari). <CARIPLO SOLARIS - Atomic Layer Deposition of metals and insulators for microelectronics and spintronics applications ( ). Objectives. The project main objective is the development of magnetic tunnel junctions (MTJ) based on the atomic layer deposition (ALD) of magnetic oxides, tunnel oxides, and magnetic metals. Parteners. CNR-INFM, Laboratorio Nazionale MDM, Agrate Brianza. Role. Coordinator (M. Fanciulli). <CARIPLO - Nonlinear optical switching of GFP protein arrays obtained by AFM surface functionalization ( ). Objectives. The project aims at developing ordered structures of fluerescent molecules and proteins on insulating, conducting and semiconducting subtrates. The fluorescence emission is a tool for visualizing the positive binding of small molecules to the surface. Moreover, the possibility of using fluorescent bistable proteins allows investigating the effects on the protein emission induced by the vicinity of insulating or conductive surfaces. The expertise gained by the proposing laboratories will be essential for the future development of optical molecular memories, based mainly on bistable fluorescent proteins. Partners. Laboratory for Advanced Bio-Spectroscopy, Dipartimento di Fisica, Università di Milano Bicocca; Laboratorio di microscopia a sonda, Dipartimento di Scienza dei Materiali, Università di Milano Bicocca; Laboratorio di Chimica Generale, Dipartimento di Chimica Generale, Università di Pavia; CNR-INFM, Laboratorio Nazionale MDM, Agrate Brianza. Role. Partner. MDM responsible: G. Tallarida. <MARTA - Metamateriali per Applicazioni Radar e Telecomunicazioni Avanzate ( ). Objectives. Metamaterials have been widely recognized as a rich class of objects able 41

43 to solve a certain number of electromagnetic applicative problems, like irradiation, propagation, and selective absorption of microwave field. The Laboratorio Nazionale MDM, in collaboration with the Consorzio Nazionale Interuniversitario per le Telecomunicazioni (CNIT), is involved in such a research on the Project Metamateriali per Applicazioni Radar e Telecomunicazioni Avanzate (MARTA). Such project focuses on innovative numerical methods, applications and materials and covers several decades of gigahertz of the microwave field. The expected results are particularly relevant in the fields of radar technology and telecommunications. Partners. CNR-INFM, Laboratorio Nazionale MDM; Consorzio Nazionale Interuniversitario delle Telecomunicazioni. Role. Coordinator (E.Prati) of Workpackage Realization and Measurements of Electromagnetic Metamaterials. European projects <NEON - Nanoparticles for Electronics ( ) (GROWTH; FP5). Objectives. The goal of this project is the controlled growth of nanocrystals buried in a thin silicon oxide layer using silicon technology compatible processes and the demonstration through this material research of the usefulness of the properties of these nanocrystals to fabricate better electronic memory devices. Fundamental issues related to the nanoclusters formation and to their characterization are among the most important objectives of the project. Partners. INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy; CNRS CEMES, Toulouse, France; National Center for Scientific Research Demokritos, IMEL, Athens, Greece; ZMD, Dresden, Germany, STMicroelectronics, Agrate, Italy; University of Aarhus, Denmark; PHASE-CNRS, Strasbourg, France; FZR, Institute for Ion Beam Physics and Materials Research, Dresden, Germany. Role. Partner and responsible (M. Fanciulli) of the Atomic scale characterization WP. Web-site: <INVEST - Integration of very high-κ dielectrics with silicon CMOS technology ( ) (IST; FP5). Objectives. The general objective is to overcome technology barriers to enhance high frequency performance and increase the density, storage capacity and functionality of ICs. The primary sub-objective is to demonstrate that complex metal oxides (CMOX) exhibiting high dielectric constant κ are suitable for the replacement of SiO 2 gate in deepsubmicron technology (<100 nm) and that they can be integrated with part of silicon CMOS technology. This will be achieved by developing appropriate methodology, based on molecular beam epitaxy (MBE), for the deposition of high-κ (k>20) CMOX materials. Our secondary sub-objective is to introduce novel materials combinations to combine increased functionality (e.g. non-volatility) and density in storage devices with enhanced performance. Partners. INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy; National Center for Scientific Research Demokritos, IMS, Athens, Greece; Interuniversitair Micro-Electronica Centrum (IMEC), Leuven, Belgium; IBM Zurich Research Lab., Switzerland; Max Planck Geselschaft zur Foerderung der Wissenschaften E.V., Halle, Germany); Oxford Applied Research Ltd., UK, TU-Clausthal, Germany, Riber S.A. Rueil-Malmaison, France; Philips Research Leuven, Belgium. Role. Partner and responsible (M. Fanciulli) of the Structural characterization WP. Web-site: <ET4US - Epitaxial Technologies for Ultimate Scaling ( ) (IST; FP6). Objectives. The goal of this project is to investigate future technology platforms, alternative to Si, for the 32 node (and beyond) devices generations. The replacements of the silicon channel with emerging materials, such as germanium and compound semiconductors 42

44 (CS), is a challenge task which will be addressed from all technologically relevant aspects: advanced large area wafers, novel gate stacks and transistor processing. The first technological objective is to demonstrate that device quality large area compliant substrates of Ge- on-insulator (GOI) and CS-on-insulator (CSOI) can be fabricated. GOI and CSOI will be grown by developing a strained oxide template on Si technology based on molecular beam epitaxy (MBE). The second technological objective is to demonstrate high quality gate stacks on Ge and CS. The challenge is to find a suitable set of high-κ materials to be used as gate dielectrics for whithout compromising the high mobility of the channel. The third technological objective is to integrate the new channel and gate materials with a 200 mm semiconductor wafer processing line to demonstrate high mobility transistors for the chosen material systems. Partners. CNR-INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy; National Center for Scientific Research Demokritos, IMS, Athens, Greece; Interuniversitair Micro- Electronica Centrum (IMEC), Leuven, Belgium; IBM Zürich Research Lab., Switzerland; Ecole Polytechnique Federale de Lausanne (EPFL), Lausanne, Switzerland; Clausthal University of Technology, Institute of Theoretical Physics, Clausthal, Germany; Philips Research Leuven, Belgium; DCA Instruments, Finland. Role. Partner and responsible (M. Fanciulli) of the WP3 Advanced gate stack materials. Web-site: <CHEMAPH - Chemical Vapor Deposition of Chalcogenide Materials For Phase-Change Memories ( ) (IST; FP6). Objectives. The project aims at the development of chalcogenides manufacturing process based on a chemical-based technique, metal-organic chemical vapor deposition (MOCVD). MOCVD enables the production of thin films with superior quality compared to those obtained by sputtering, especially in terms of conformality, coverage, and stoichiometry control, and allows implementation of phase-change films in nanoelectronic devices. The main phase-change chalcogenide material system that will be investigated is Ge 2 Sb 2 Te 5 (GST), as it is already the basis of optical storage media and prototype PCM devices. This objective will require extensive studies of the thermochemical properties of a variety of possible precursor materials, and their interactions, and the investigation of the optimal process conditions to achieve desired film properties, such as resistivity, phase-transition temperature, roughness, density. The optimized deposition process will be applied for fabricating state-of-the-art electrical memory cells at the 90/65nm node. The device performance will be evaluated using standard sputter-deposited devices as a benchmark, in terms of parameters such as the programming current, cyclability, and retention. Partners. Laboratorio Nazionale MDM CNR-INFM, Agrate Brianza; Italy, Aixtron AG, Germany, ST Microelectronics, Agrate Brianza, Italy; Epichem Limited, United Kingdom; Consejo Superior De Investigaciones Cientificas, Madrid, Spain; Vilnius University, Lithuania. Role. Coordinator (Claudia Wiemer). Web-site: <REALISE - Rare earth oxide atomic layer deposition for innovation in electronics ( ) (IST; FP6) Objectives. The project aims at (a) deposit high permittivity rare earth oxide layers with subnanometer control and (b) integrate these into innovative memory and communications devices. The process that is the subject of this project is atomic layer deposition (ALD), the leading technology for deposition of nanometre-scale films. Rare earth oxides show promise as high-permittivity dielectrics in nano-electronic devices, but no satisfactory ALD process exists for these oxides. Device performance is currently limited by the incorporation of impurities in the material, morphological instability and a poor interface to silicon. The project aims to overcome these difficulties through project goals that span the entire ALD process: (i) design, synthesis and scale-up of suitable precursors, (ii) high-resolution characterisation of film quality and interface to semiconductor (Si, Ge), 43

45 and (iii) optimisation of deposition parameters. In order to transform this ALD process into industrial production and demonstrate its use as an enabling technology for high-volume nano-electronics, REALISE will also (i) scale-up new ALD process to industrially-sized Si wafers, and (ii) integrate dielectric layers into test capacitors for innovative memory (DRAM), embedded electronics (NVM) and wireless (RF) applications. In the longer term, this technology will enable down-scaling of mass-produced complementary metal-oxidesemiconductor (CMOS) devices. Partners. Tyndall National Institute, Cork, Ireland; University of Liverpool, UK; Epichem, Bromborough, UK; University of Helsinki, Finland; ASM-Microchemistry, Helsinki, Finland; CNR-INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy; CEMES-CNRS, Toulouse, France; Qimonda, Dresden, Germany; NXP, Eidhoven, The Netherlands; STMicroelectronics, Agrate Brianza, Italy. Role. Partner and responsible (M. Fanciulli) of WP3 Interface engineering for NVM and CMOS. Web-site: <VERSATILE - Vertically stacked memory cells based on heterojunctions made of hybrid organic/inorganic materials ( ) (IST; FP6). Objectives. The project aims at developing non-volatile memories based on cross-bar architectures. In particular we will integrate junction made of II-VI and organic/polymeric semiconductors into cross-bar type memories to obtain a scalable cross-bar non volatile memory with both bit storage and selection elements vertically stacked one on top of the other. The junction material should require a low thermal budget for its preparation to be compatible with the most promising technologies for the bit storage element, such as chalcogenides. Partners. CNR-INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy; STMicroelectronics, Agrate Brianza, Italy; Institute for Nanoelectronics Technical University of Munich, Germany;The Danish Polymer Centre, Denmark; Risoe National laboratory Institute of Physics, Polish Academy of Sciences, Poland. Role. Coordinator (S. Ferrari) Web-site: <EMMA - Emerging Materials for Mass-storage Architectures ( ) (IST; FP6). Objectives. This project will investigate the feasibility of emerging new non-volatile memory concepts based on resistive-switching materials for enabling new mass-storage memory systems. These new memory concepts allow integration of the memory element in contact and interconnect structures resulting in very small memory cells and even offer the possibility of 3-D memory layer stacking. These new memory solutions are needed for the sub-32 nm integration technology nodes where current memory concepts will no longer scale. The program will study high-density resistive switching non-volatile memories, including binary resistive switching oxides and Cu-TCNQ. Focus will be on concept scalability, based on gained understanding of the physical operation concepts. Investigation will further include cell integration aspects, reliability assessment, and memory architectures. Partners. Interuniversitair Micro-Electronica Centrum (IMEC), Leuven, Belgium; CNR- INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy; STMicroelectronics, Agrate Brianza, Italy; IUNET- Consorzio Nazionale Interuniversitario per la Nanoelettronica, Italy; RWTH, Aachen, Germany; CNRS-L2MP, France. Role. Partner and responsible (M. Fanciulli) of the WP1 Memory materials and electrode technology Web-site: <FOREMOST - Integration of Fourty Five Nanometers CMOS Technology ( ) (Medea+ 2T103). Objectives. The project aims to develop advanced process modules and transistors architectures to realize the demonstration of a full CMOS 45 nm process technology 44

46 in both European manufacturing industrial facilities of Crolles 2 Alliance (STMicroelectronics, Philips Semiconductors, Freescale Semiconductor) and Infineon Technologies. For CMOS logic process, the first validation of the integration on a complex test vehicle is proposed in Q4 07 and the reliability of this process for Q2 08, while the 45 nm demonstrator is scheduled for Q1 08 and the test and analysis of DSM effects in Q2 08. Specific DRAM process steps will also be available then. Partners. STMicroelectronics, Crolles, France and STMicroelectronics, Agrate Brianza, Italy; Air Liquide, France; Aixtron AG; ASM international,, Carl Zeiss STM, Germany; CEA LETI, France; Fraunhofer Center for Nanoelectronic Technologies, Germany; Epichem Limited, United Kingdom; Freescale, IBS, France; Institute of Microelectronics, Electromagnetism and Photonics, France; Institute for Microelectronics Stuttgart, Germany; Jordan Valley, Laboratory of Hyperfrenquencies and Characterization of University of Savoie, France; Leica Microsystems, Laboratoire des Matériaux et du Génie Physiquem CNRS- INPG, France; LTM/CNRS-UJF, France; CNR-INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy; NCSR Demokritos, Greece; Philips Semiconductors, Crolles, France. Role. Partecipant in WP2 Gate stack modules. MDM responsible: M. Fanciulli. <TCAMMD - Thermal Characterization of Advanced Material in Ultra-scaled Microelectronic Devices (August 2006-July 2007) (IEF Marie Curie; FP6). Objectives. The goal of this project is to implement thermal characterization techniques capable of addressing the device nano-scale. To achieve this objective the experimental technique is combined with the heat transfer modelling in the device under investigation. The project will focus on specific microelectronic devices such as ultra-scaled metal-oxidesemiconductor field-effect transistors (MOSFETs) as well as novel non-volatile memories (NVM) based on phase-change materials (PCM). In the latter case the determination of the thermal properties of the active material (calchogenides) and of the heater (TiN) is very important for the device functional characterization. An additional objective of the project is also to use the developed technique to address the thermophysical properties of thin films, such as oxides with high and low dielectric constant, which today are considered as substitutes of silicon oxide both in the front as well as in the back end of near future CMOS technology. Partners. Laboratoire TREFLE, Université Bordeaux 1, France ; CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy; STMicroelectronics, Agrate Brianza, Italy. Role. Host Institution. Other International Projects <MAE - Italia-Russia joint research project in science and technology Topic: Material Science. Novel high dielectric constant films for CMOS structures (2004). Objectives. The main objectives of this project are the synthesis of Lu based precursors and the growth by ALD of Yb 2 and Lu 2 films, their structural, physical, chemical, electrical, electronic, and interface characterization. The interest in Yb 2 and Lu 2 films has its origin in their potential applications as gate and interpoly dielectrics in ultra scaled CMOS devices. Within the project, ALD growth parameters will be established and optimized, given that some commercially available potential precursors exist. Moreover, at least two novel Lu based precursors will be synthesized for the ALD growth of Lu 2. Such precursors have not been produced before. This work is relevant given that the precursor synthesis and the optimization growth parameters are among the main research fields related to ALD. The optimization of the growth parameters will be based on the films morphological, structural, chemical, electrical, defective and interfacial properties. Partners. INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy; Physics of Solids Department of the Moscow Engineering Physics Institute, Russia; Technology of 45

47 Organometallic Compounds Department of the Institute of Organometallic Chemistry, Nizhny Novgorod, Russia. Role. Coordinator (M. Fanciulli). <MAE - Italia-Austria joint research project in science and technology Topic: Material science. Development of a comprehensive Scanning Probe Microscopy tool for the characterization of future high-κ dielectric thin films ( ). Objectives. Goal of this collaboration is to set-up an AFM based measurement extension to allow the local characterisation of electrical properties of high-κ dielectrics. The expertise developed by the tow laboratories will be combined to build up a module and a reliable measurement procedure for a comprehensive characterisation of local electrical properties of high-κ dielectrics. KPFM and conductive-afm methods will be integrated for the measurement of flat band voltage and of leakage current through the probe/dielectric/semiconductor stack. The purpose is to discriminate the mechanisms responsible for the conduction through the dielectric layer and to connect them with film properties, film growth conditions and post- deposition treatments. Partners. CNR-INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy; Department of Physics, University of Leoben, Austria. Role: Coordinator of the Italian unit (G.Tallarida). <MAE - Italia-Poland joint research project in science and technology - Topic: Material Science Experimental and theoretical investigation of the interface between high-κ dielectrics and semiconducting substrates (2005). Objective. The objective of the project is to grow, by atomic layer deposition, and characterize thin high-κ dielectric layers on Si, SiGe, Ge, an also III-Vs (GaAs, GaN). High-κ dielectrics on semiconductors may lead to increase functionalities of conventional Si-based CMOS technology as well as to the realization of novel devices using other group IV or group III-V materials. The applications in microelectronics, nanoelectronics, and spintronics depend critically on the semiconductor/oxide interface which will be the main topic addressed during the first year of the planned activity. Partners. CNR-INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy; Institute of Physics of the Polish Academy of Science, Poland. Role. Coordinator (M. Fanciulli). <VAMAS Project A10: X-ray reflectivity measurements for evaluation of thin films and multilayer thickness ( ). Objectives. The Versailles Project on Advanced Materials and Standards (VAMAS) supports trade in high technology products through international collaborative projects aimed at providing the technical basis for drafting codes of practice and specifications for advanced materials. VAMAS activity emphasizes collaboration on pre-standards measurement research, intercomparison of test results, and consolidation of existing views on priorities for standardization action. Through this activity, VAMAS fosters the development of internationally acceptable standards for advanced materials by the various existing standards agencies. Project A10 will assess the accuracy and precision of thickness, density, and roughness measurements of thin-film structures determined by grazing-incidence X-ray reflectance (XRR) in an interlaboratory comparison. This project is believed to be timely in view of the growing use of XRR for characterizing thin-film materials and the recently completed pilot study of measurements of the thicknesses of 1.5 nm to 8 nm SiO 2 films on silicon. This pilot study was conducted under the auspices of the Consultative Committee on Amount of Substance and involved thickness measurements by SIMS, XPS, XRR, and seven other methods. Partners. University of Brescia, Italy; Bruker AXS, Germany; University of Tsukuba, Japan; Masaryk University, Czech Republic; University of Bratislava, Slovak Republic; CNR- INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy; IMM-CNR, Bologna, Italy; 46

48 Univeriste de Montpellier II, France; Université du Maine, France; University of Durham, United Kingdom; Pilkington European Technical Centre, United Kingdom; Bede plc., United Kingdom; Physikalish-Technische Bundesanstalt, Germany; National Institute of Advanced Industrial Science and Technology Tsukuba, Japan; Material Analysis & Test team, RIST, Korea; Institute of microelectornics technologies RAS, Russia; Jordan Valley Semiconductor, USA, University of Albany, USA. Role. Partner. MDM responsible:c. Wiemer. Web-site: 47

49 International Collaborations ø Aixtron, Aachen, Germany ø Brazzaville University, Congo ø Bruker BioSpin, Rheinstetten, Germany ø CEA, Tours, Tours University, France ø CERN ISOLDE, Geneva, Switzerland ø CNRS CEMES, Toulouse, France ø CNRS PHASE, Strasbourg, France ø CSIC, Madrid, Spain ø Department of Physics, Boston University, USA ø Department of Physics and Astronomy, University of Leeds, UK ø Ecole Nationale Supérieure d Arts et Métiers, France ø Epichem, Liverpool, United Kingdom ø ETH, Swiss Federal Institute of Technology Zurich, Switzerland ø FZR, Institute for Ion Beam Physics and Materials Research, Dresden, Germany ø Interuniversitair Micro-Electronica Centrum (IMEC), Leuven, Belgium ø IBM Zurich Research Lab, Zurich, Switzerland ø IHP - Frankfurt am Oder, Germany ø IMEC, Belgium ø INFM-OGG c/o ESRF, GILDA CRG, Grenoble, France ø Institute of Theoretical Physics, EPFL, Lausanne, Switzerlad ø Institute of Physics and Astronomy, University of Aarhus, Denmark ø Institute of Physics, University of Leoben, Austria ø Institute of Organometallic Chemistry of the Russian Academy of Sciences, Nizhny Novgorod, Russia ø Institute of Physics, Polish Academy of Sciences, Warsaw, Poland ø IPCMS-CNRS, Strasbourg, France ø Max Planck Institute of Microstructure Physics, Halle (Saale), Germany ø Max Planck Institute for Biochemistry, Department of Membrane and Neurophysics, Martinsried, München, Germany ø MEPhI, Moscow Engineering Physics Institute, Russia ø National Center for Scientific Research DEMOKRITOS : IMS and IMEL, Athens, Greece ø Oxford Applied Research Ltd, Oxon, UK ø Philips Research Leuven, Belgium ø Riber S.A., Rueil-Malmaison, France ø STMicrolectronics, Crolles, France ø Technical University of Muenchen - Institute for Nanoelectronics, Germany ø The Danish Polymer Centre - Risø National Laboratory, Denmark ø TU-Clausthal, Germany ø Tyndall Institute, Cork, Ireland ø University of Bordeaux, Bordeaux, France ø University of Florida, USA ø Vilnius University, Vilnius, Lithuania ø Walter Schottky Insitut, Am Coulombwall Garching, München, Germany ø ZMD, Dresden, Germany 48

50 National Collaborations ø Consorzio Nazionale Interuniversitario delle Telecomunicazioni, Firenze ø CRS-Democritos, CNR-INFM Trieste ø Dipartimento di Chimica Generale, Università di Pavia ø Dipartimento di Fisica, Università di Bologna ø Dipartimento di Fisica, Università di Brescia ø Dipartimento di Fisica, Università di Cagliari ø Dipartimento di Fisica, Università di Catania ø Dipartimento di Elettronica e Informazione e Dipartimento di Fisica, Politecnico di Milano ø Dipartimento di Fisica e Dipartimento di Scienze dei Materiali, Università di Milano Bicocca ø Dipartimento di Fisica, Università di Modena ø Dipartimento di Fisica, Scuola Normale Superiore, Pisa ø Dipartimento di Ingegneria ed Elaborazione dell Informazione, Università di Pisa ø Dipartimento di Ingegneria Elettronica, Università La Sapienza di Roma ø Dipartimento di Fisica, Politecnico di Torino ø Dipartimento di Ingegneria dei Materiali e Dipartimento di Fisica, Università di Trento ø Dipartimento di Fisica, Università degli studi di Trieste ø Istituto di Scienze e Tecnologie Molecolari, CNR Milano ø ItalStructures s.a.s., Riva del Garda, Trento ø Laboratorio Nazionale INFM-TASC, Trieste ø STMicroelectronics, Agrate Brianza 49

51 1. Scaling issues in CMOS logic and memory devices 1.1 Structural and electrical properties of Hf-based oxides for interpoly applications 1.2 Al 2 as interpoly dielectrics in non-volatile memory devices 1.3 HfO 2 as gate dielectrics for ultra-scaled CMOS devices 1.4 Nanoscale electrical properties of HfO 2 and ZrO 2 thin films studied by conducting atomic-force microscopy 1.5 Thermal stability of HfO 2 /TiN gate stacks for 45 nm CMOS devices 1.6 Oxygen diffusion in HfO 2 /SiO 2 /Si stacks 1.7 Characterization of the mechanical stress induced in silicon during device fabrication 1.8 Low-k materials for intra-metal dielectrics 1.9 Advanced materials for interconnects According to the International Technology Roadmap for Semiconductors (ITRS - the main trends in the semiconductor industry for the next 15 years will follow two principal routes. Complementary metaloxide-semiconductor (CMOS) devices will remain the fundamental building block at least up to the year 2020, undergoing continuous shrinking and architectural changes. On the other hand, new devices will be introduced in the latter half of the next decade utilizing different and new ways of processing and storing information. Scaling planar bulk CMOS introduces significant challenges, related not only to technological issues, but also to fundamental topics in material science and characterization techniques. In this chapter, we describe the efforts of MDM in addressing some of these issues. One of the most difficult challenges is the reduction of the equivalent oxide thickness (EOT), required for allowing device density increase, whilst keeping suitable drain current. The currently established solution is to replace conventional SiO 2 gate oxides with a material having higher permittivity (κ). High-κ insulators can be grown physically thicker for the same (or thinner) EOT, thus offering significant gate leakage reduction. With the substitution of SiO 2, also the polycrystalline silicon gate electrode becomes obsolete, and new materials, with characteristics tailored on the selected high-κ insulator, are to be developed. In Flash memory devices, on the other hand, continuous scaling and the reduction in write voltage requires the use of a thinner interpoly and tunnel oxide. To assure retention and ease of erase/write in tunnel oxides, as well as constant coupling ratio in interpoly dielectrics, high-κ material with specific properties are to be introduced into Flash memory process, as well. In this field, MDM has built a solid and extensive expertise, ranging from the deposition and processing of Hf- and Albased materials, to their structural and electrical characterization. The first six 50

52 paragraphs of this chapter highlight this effort. Hf-based oxides and Al 2 are proposed as interpoly dielectrics (paragraph 1.1 and 1.2). The structural and electrical characterization of these materials addresses some topics, such as thermal stability and dielectric quality, relevant in view of their integration in non-volatile memory devices. HfO 2 has been widely investigated as possible gate dielectrics. However, several issues mainly related to the electrical properties of the film and of its interface with silicon, as well as the definition of the optimal deposition process, remain open. These are addressed in paragraph 1.3 that describes the structural and electrical properties of hafnium oxide films obtained by atomic layer deposition, and in addition focuses on the conduction mechanisms responsible for the leakage current. The spatial distribution of leakage current and how it correlates to the structural properties of ALD grown HfO 2 and ZrO 2, is investigated by conducting-afm and is the subject of paragraph 1.4. In paragraph 1.5 the characterization of structural and chemical properties of HfO 2 /TiN stacks is reported, evidencing their influence on the measured EOT. Finally, paragraph 1.6 focuses on the mechanisms responsible for the oxygen exchange at the HfO 2 /SiO 2 /Si interfaces. The mechanical stress induced in silicon during device fabrication has a remarkable impact on the device performances and reliability, and its characterization and control is mandatory to ensure proper device functionality. On the other hand, the controlled introduction of strain in the silicon channel region is being investigated as a method for enhancing carrier mobility. However, the continuous device scaling makes the conventional characterization techniques and the predictive modelling tools increasingly inadequate. On this topic, MDM has been successful in using μraman spectroscopy for characterising the stress induced into silicon during device fabrication, and by exploiting resonant Raman spectroscopy for the characterization of strain in arrays of subwavelength devices. The main achieved results in this field are reported in paragraph 1.7. Also the so called back-end processes introduce interesting topics in material science. Indeed, to minimize signal propagation delay and power consumption, development of low dielectric constant (lowκ) material together with low-resistivity metal system is critical. Low-κ material should have sufficient mechanical, chemical, and thermal integrity to survive severe integration processes, such as chemical mechanical planarization, etching/ashing/ wet cleaning, and assembly/packaging. On the other hand, resistivity of narrow Cu interconnects is bound to increase as the line width shrinks, due to the increasing impact of electron scattering at the Cu/barrier interface and the grain boundary. Barrier engineering including construction of very thin and low-resistive barrier metal, as well as stable interface with low-κ material, is essential to achieve high conductivity in narrow Cu interconnects. Paragraphs 1.8 and 1.9 address these issues: in the first, two porous silicate (SiOCH) films are compared, focusing on their structural properties and their stability against two different etching cleaning steps; finally, the Cu/barrier system is thoroughly investigated, and a few amongst the most suitable barrier materials and deposition processes for very narrow interconnects are explored. The activities presented in this chapter are carried out in close collaboration with STMicroelectronics, through the industrial project renewed every year, as well as in the framework of several research projects. The latter include the EU FP5 IST project INVEST ( ), focused on the search for high-κ dielectrics alternative to SiO 2, the EU MEDEA + project FOREMOST ( ) aiming at the development of advanced process modules and transistor architectures for a full CMOS 45nm process technology in Europe, the National FIRB Project Miniaturized systems for electronics and photonics RBNE012N3X ( ), funded by the Italian Ministry of Research, and the Italia-Austra joint research project in science and technology ( ), supported by the Ministry of Foreign Affairs, devoted to the development of a scanning probe microscopy tool for the characterization of high-κ dielectric thin films. 51

53 1.1 - Structural and electrical properties of Hf-based oxides for interpoly applications C. Wiemer 1, R. Piagge 2, S. Spiga 1, E. Bonera 1, M. Fanciulli 1, M. Alessandri 2, G. Ghidini 2, M. Caniatti 2, A. Sebastiani 2, D. Caputo 2 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy 12 STMicroelectronics, Agrate Brianza, Italy Due to its high dielectric constant (κ) and relatively high conduction band offset with silicon, HfO 2 is currently considered to substitute SiO 2 as gate dielectric in logic devices (see paragraphs ). For interpoly dielectrics, where the requirements for leakage current are more severe than for logic devices, Al 2, HfO 2 /Al 2 nano-laminates with different thickness and composition, and hafnium-based materials (such as hafnium aluminates and silicates) are evaluated [1]. Besides the low leakage current, due to the high thermal budget involved during device fabrication, the interpoly dielectrics has to demonstrate good thermal stability; moreover, the evolution with temperature of the dielectric and the structural properties have to be thoroughly investigated. Especially this last point is addressed in this paragraph, focused on Hf aluminates and silicates, and nanolaminates. Hf 1-x Al x O y films formed by 26% of Al 2 and 74% of HfO 2 were grown by ALCVD TM at 300 C on RCA precleaned p-type silicon wafers. Al(CH 3 ) 3, HfCl 4 and H 2 O have been used as precursors. Post-deposition annealing (PDA) was carried out in ultra pure N 2 at atmospheric pressure at 900 C for 2 and 30. Chemical composition and thickness uniformity have been checked by X-ray fluorescence spectroscopy and ellipsometry, respectively. No changes in the Hf/Al/O content after post deposition annealing and the same composition for different thickness have been demonstrated by X-ray fluorescence characterization. However, the increase of the film electronic density and the decrease of the thickness measured by X-ray reflectivity after 2 minutes of annealing evidence the densification of the high-κ layer. After 30 minutes of annealing, the thickness of the high-κ is not found to vary, while the variations in the electronic density depend on the film thickness: the electronic density decreases for thinner films, does not change substantially for 10 nm thick films, and increases for thicker films. These modifications suggest the occurrence of some chemical variation due to the long annealing time, namely diffusion of Si from the substrate, which might more strongly affect the electronic density of very thin films. The thickness of the interlayer (IL) between the high-κ layer and Si increases with annealing. Its electronic density value also increases. After 2 minutes of annealing, the electronic density is still comparable to that of SiO 2, whereas after 30 minutes of annealing it is even higher than that of crystallized SiO 2. While a small increase of the IL electronic density can be due to a non gaussian-distributed roughness at the high-κ/il interface, as it can be the case at an amorphous/ crystallised interface, a strong increase can be due to a real intermixing between the high-κ and the IL and/or to some diffusion from the substrate. Therefore, after 30 minutes of annealing, not only the high-κ, but also the IL, experiences some chemical variations. All the as-deposited samples are amorphous, whereas all the annealed films are crystallised in the orthorhombic Hf 1-x Al x O 2 phase. No variations are found in the diffractogram of films annealed for 2 or 30 minutes. The degree of crystallization does not depend on the film thickness, as shown in Figure 1, i.e. if an amorphous component is present, its percentage does not depend on the film thickness. The lattice volume is higher in thinner films, while grain size and lattice strain are higher in thicker films [2]. Figure 1. XRD data of Hf 1-x Al x O 2 films with different thickness (black lines). The simulations (red lines) are obtained starting form the 8.1 nm thick sample and varying only thickness parameter between the samples. 52

54 different layers is clearly depicted in Figure 2 where the X-ray reflectivity curves before and after annealing of nano-laminates with a basic sequence formed by 2nm Al 2 and 1nm of HfO 2, are presented. Figure 2. XRR data of HfO 2 /Al 2 stacks formed by sequences of 1 nm of HfO 2 and 2 nm of Al 2 before (black) and after (blue) PDA. Inset: corresponding electronic density profiles. The equivalent oxide thickness (EOT) and dielectric constant of Hf 1-x Al x O y films formed by 26% of Al 2 and 74% of HfO 2 are measured by capacitancevoltage characteristics of capacitors with Al and Au metal gates. The extraction of the dielectric constant was obtained from the slope of the linear fit of the EOT data versus the Hf 1-x Al x O 2 films thickness. The intercept of the linear fit with the y-axis gives the EOT of the interfacial layer. EOT values are extracted without taking into account quantum mechanical corrections. The dielectric constant of as deposited Hf 1-x Al x O y films was determined to be 16(±2). An increase of the dielectric constant is measured after annealing at 900 C for 2 minutes (Table I). Also a longer thermal treatment is effective in increasing the dielectric constant of the as deposited Hf 1-x Al x O y films [3]. The dielectric constant behaviour of nano-laminate materials is driven by the intermixing phenomenon and by the crystallization, as well as by the relative Al 2 / HfO 2 concentration, with lower κ being measured on Al 2 -rich stacks. For HfO 2 -rich nano-laminates, the highest κ value (19) is obtained when crystallization in the orthorhombic phase occurs. When the physical thickness of the hafnia layers is above 4 nm there is no intermixing after annealing, the alumina-hafnia layers are still well separated and the hafnia clearly crystallizes in the monoclinic phase. In this case, the κ value (17) does not appreciably change after annealing [1]. Hafnium silicates deposited by MOCVD with two compositions (40% and 60% SiO 2 ) and two nitridation conditions (low nitrogen content: DPN1, high nitrogen content: DPN2) have been evaluated [5]. XRD analysis reveals that after DPN films with higher N content are less crystallized than films with lower N content. Additional annealing at 1000 C (PDA) does not have any impact on the crystallization for both high- and low-n content films. In contrast with what assessed for hafnium aluminates, clear evidence of phase separation is found with annealing of hafnium silicates. In FT-IR spectra (Figure 3), the shift of the absorption band from 1010 to 1080 cm -1 after PDA can be ascribed to a modification of the environment Hf 1-x Al x O 2 κ EOT IL (nm) As grown 16.0±2 1.40± C, 2 min 22.7±2 1.20± C, 30 min 24.2±1 1.50±0.15 Table I. Hf 1-x Al x O 2 dielectric constant (κ) and EOT of the interfacial layer (IL). HfO 2 /Al 2 nano-laminates with different relative thickness and compositions have been deposited by ALD at 300 C. Each sample is formed by four Al 2 / HfO 2 sequences, and by a final Al 2 layer, so that the structure of each sample is: Si/SiO 2 /(Al 2 /HfO 2 )x4/ Al 2. Combined transmission electron microscopy and X-ray reflectivity analysis show that when the physical thickness of the HfO 2 layers is below 4 nm, an intermixing of alumina-hafnia layers occurs after PDA. The onset of crystallization after annealing depends on the relative HfO 2 content [4]. The intermixing of the Figure 3. Absorbance spectra for as-grown and annealed hafnium silicates. 53

55 of the Si-O atoms. While in the as-grown films the position of the absorption band is influenced by the Hf atoms, this is not the case for the annealed films, where the position of the absorption band corresponds to the one reported in literature for SiO 2. In the amorphous as grown films, FTIR analysis evidences that these films are homogeneous silicates, with Hf atoms in the neighbourhood of Si-O atoms. All the Hf-rich films are crystallized after PDA in the orthorhombic phase of HfO 2. The crystallization of Si-rich films is detected by TEM. In the as-grown samples, the lower frequency of the infrared band could be explained by the involvement of some heavy Hf ions in the vibration. In the annealed films, the fact that the absorption band is centred at the value of bulk silica suggests that segregation of SiO 2 is taking place. We can therefore argue that this annealing results in the separation of the homogeneous silicate into a Hf-rich crystallized silicate and a Si-rich amorphous silicate. κ value increases with decreasing the SiO 2 content, as expected. In samples with 60% SiO 2, the dielectric constant value is not influenced by the N content, while the 40% SiO 2 samples show a κ dependence on N content: increasing the N content a decrease of the κ value is achieved on the as deposited samples, while a κ increase is observed after 1000 o C annealing on the high N content films [1]. The κ value always remains lower than 12 and lower than the one measured for hafnium aluminates with nominally the same Hf content. [1] M. Alessandri, R. Piagge, S. Alberici, E. Bellandi, M. Caniatti, G. Ghidini, A. Modelli, G. Pavia, E. Ravizza, A. Sebastiani, C. Wiemer, S. Spiga, M. Fanciulli, E. Cadelano, G.M. Lopez, V. Fiorentini, ECS Transactions 1(5) (2006). [2] C. Wiemer, E. Bonera, R. Piagge, Structural and morphological characterization of Hafnium Aluminates deposited and annealed by ST, MDM-D [3] S. Spiga, R. Piagge and D. Caputo, Electrical characterization of HfO 2 and HfAlO films deposited by ALD using the ASM Pulsar 3000 reactor, MDM-D [4] C. Wiemer, R. Piagge, Structural and morphological characterization of HfO 2 /Al 2 stacks deposited and annealed by ST, MDM-D [5] C. Wiemer, E. Bonera, R. Piagge, Structural and morphological characterization of demo samples of hafnium silicates, MDM-D Al 2 as interpoly dielectric in non-volatile memory devices C. Wiemer 1, R. Piagge 2, S. Spiga 1, E. Bonera 1, V. Fiorentini 3, M. Fanciulli 1, M. Alessandri 2, G. Ghidini 2, A. Del Vitto 2, M. Caniatti 2, A. Sebastiani 2, S. Alberici 2, G. Pavia CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy STMicroelectronics, Agrate Brianza, Italy SLACS-INFM, Departmento di Fisica, Università di Cagliari, Italy Based on a preliminary screening of the electrical performances of capacitors incorporating high-κ dielectrics, Al 2 was selected as the most promising candidate to substitute the current oxide-nitride-oxide (ONO) interpoly dielectric [1], at least for the incoming technology nodes, where a moderately high-κ is sufficient. Amorphous Al 2 films grown by atomic layer deposition (ALD) at 300 C from TMA and water are considered. The structural evolution of films in the 5-60 nm thickness range is studied for annealing temperatures up to 1100 C. Annealing causes shrinking and densification of the Al 2 layer, whereas its interface with silicon remains unchanged. The XRR data and simulation of as-grown and annealed Al 2 films are shown in Figure 1. After annealing at 900 C in N 2, the films are partially crystallised, with amorphous regions and small grains with inter-planar spacing corresponding to those of several Al 2 polymorphs. After annealing at higher temperatures, the system stabilises in the γ phase of Al 2, with grain size smaller than the film thickness and with a strong preferential orientation. According to XRD in Bragg-Brentano configuration the majority of the Al 2 crystallites is found to be actually oriented with the {111} planes parallel to the Si substrate [2]. The dielectric tensor of γ alumina is derived from first-principles gradientcorrected density-functional theory calculations. The theoretical value averaged for randomly-oriented polycrystalline γ-alumina is 11.6±10%. The zz component of the dielectric tensor is 8.9: this is therefore the expected value for metal-al 2 -Si structures with alumina (111) planes parallel to the Si surface. The calculated infrared spectrum is 54

56 The dielectric constant of as grown and annealed Al 2, as well as the CET (Capacitance equivalent oxide thickness) of the interlayer developing between the dielectric and the Si substrate, are evaluated from C-V characteristics of metal-al 2 -semiconductor structures, with the Al 2 thickness varying in the 5-30 nm range. The dielectric constant of the as grown Al 2 is 8.3±0.3, while the CET of the interlayer is 0.9±0.2 nm (Figure 3). Annealing at 1030 C causes an increase of the accumulation capacitance with respect to the as deposited films. In addition, the annealing at 1030 C for 30 s in N 2 leads to an increase of the interface trap density. The extracted dielectric constant Figure 1. XRR analysis of nominally 15 nm thick Al 2. Data (blue circles: as grown, red circles: annealed) and corresponding simulations (lines). The extracted thickness are 14.2 nm and 12.2 nm for the as-grown and the annealed sample, respectively. Figure 3. CET values as a function of the Al 2 thickness for the as-deposited (black circles) and annealed samples (red squares). Gate area: (7.7±0.3) x 10-4 cm 2. Figure 2. (a) Calculated spectrum from first principles gradientcorrected density-functional theory calculations. (b) Experimental spectrum and simulation from transmission infrared spectroscopy. shown in Figure 2(a), where the frequencies are reported with their oscillator strength. The experimentally measured infrared response is reported in Figure 2(b), showing different transverse optical phonon absorption bands located at 310, 365, 515, 730, and 810 cm -1. The comparison of the two panels reveals the agreement between the calculated and measured frequencies. Qualitatively, also the calculated and measured distribution of oscillator strengths correspond. of crystalline Al 2 (Figure 3) is 9.3±0.4, in excellent agreement with the theoretical estimation, while the CET of the inter-layer is 0.8±0.2 nm [3]. Al 2 was therefore integrated in proper interpoly device structures for further electrical characterization. The test structure is based on a capacitor with a doped poly-silicon floating gate on which the dielectric material to be evaluated is deposited. A tunnel oxide isolates the floating gate from the substrate and the control gate is also made of doped poly-silicon. A method based on C-V shifts after a prolonged stress at low voltages has been setup in order to evaluate very low leakage currents, in the range of interest for the Flash memory applications. A C-V measurement is first made on the virgin structure and then after a stress at low positive voltage applied for a long time. The flat band voltage shift detected can be used to calculate the leakage current achieved during the stress, due to charge (electron) extraction from the floating gate and moving toward the positive control gate. As compared to standard 14.5 nm EOT ONO-based 55

57 capacitors, where the CV shift is not detectable, the C-V shift is very high for Al 2, 8 nm EOT, after PDA at 900 C. The shift is concentrated in the first 4 hours of stress; after that it is not detectable. A major improvement is achieved with PDA at 1030 C. No C-V shift is detected on capacitor based on Al 2 annealed at 1030 C. At 1030 C, all the Al 2 is crystallized, while it is only partially crystallized at 900 C. The calculated leakage current after 1030 C annealing (4x10-13 A/cm 2 ) is close to the required target, which is 1x10-13 A/cm 2 at a field of 4MV/cm. Recent analyses have also addressed the impact of post deposition treatments in O 2 and in NH 3 on the structural and electrical properties of Al 2 [4]. Annealing in NH 3 promotes nitrogen incorporation in the dielectric and reduces the crystallization after PDA. Annealing in O 2 favours oxygen incorporation and prevents nitrogen diffusion during PDA, while maintaining the same degree of crystallization reached after only PDA in N 2 [5]. On one hand, treatments in NH 3 or O 2 before PDA do not induce remarkable dielectric constant variation in the Al 2 layer. On the other hand, preliminary analysis on text structures have shown that these treatments might have an impact on the C-V shift after electrical stress, therefore demonstrating a possible way to further improve the leakage current of Al 2 for interpoly applications. [1] M. Alessandri, R. Piagge, S. Alberici, E. Bellandi, M. Caniatti, G. Ghidini, A. Modelli, G. Pavia, E. Ravizza, A. Sebastiani, C. Wiemer, S. Spiga, M. Fanciulli, E. Cadelano, G.M. Lopez, V. Fiorentini, ECS Transactions 1(5), (2006). [2] C. Wiemer, S. Spiga, E. Bonera, M. Fanciulli, R. Piagge, M. Alessandri, M. Caniatti, G. Pavia, E. Cadelano, G. M. Lopez, V. Fiorentini, Structural and electrical properties of annealed Al 2 thin films as inter-poly dielectric for Flash Memory applications, WoDiM (workshop on dielectrics in microelectronics) Santa Tecla (CT, Italy) June 26 th -28 th [3] S. Spiga, C. Wiemer, R. Piagge, D. Caputo, Electrical characterization of Al 2 films: as deposited and annealed films at 1030 C, N 2, 15. Determination of the dielectric constant of amorphous and crystalline films, MDM-D [4] R. Piagge, M. Caniatti, A. Del Vitto, C. Wiemer, G. Pavia, S. Alberici, E. Bellandi, A. Nale, and M. Alessandri, accepted for pubblication in ECS Transaction [5] C. Wiemer, R. Piagge, Characterization of Al 2 exposed to thermal treatments in O 2 and NH 3, MDM-D HfO 2 as gate dielectric for ultra-scaled CMOS devices S. Spiga 1, C. Wiemer 1, G. Scarel 1, S. Baldovino 1, G. Tallarida 1, M. Fanciulli 1, C. M. Compagnoni 2, A. S. Spinelli 2, A. Bianchini 2, A. L. Lacaita 2 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy 12 Dipartimento di Elettronica e Informazione, Politecnico di Milano-IU.NET, Milano, Italy HfO 2 is recognized as a good candidate to replace SiO 2 in future microelectronics technology nodes [1,2], due to its high dielectric constant (16-25), large band gap ( ev) and conduction band offset (2 ev) [3], equivalent oxide thickness (EOT) down to nm with low leakage currents, low density of interface states and good electrical reliability [4,5]. On the other hand, HfO 2 suffers from large transient currents, which may affect the functionality of several integrated circuits (ICs) [6]. In this work we characterized the structural/ chemical and electrical properties of HfO 2 films and HfO 2 /Si interface, as well as transient currents in HfO 2 dielectrics. HfO 2 films were grown by atomic layer deposition (ALD) on chemically oxidized or H-terminated Si(100) at various temperatures ( C) and using various precursor combinations. H 2 O,, Hf(O t Bu) 2 (mmp) 2, Hf(mmp) 4 were used as oxygen precursor, while HfCl 4, Hf(O t Bu) 2 (mmp) 2, Hf(mmp) 4 as metal precursor. The precursors choice, as well as film thickness, growth temperature and post-deposition annealing, influence the film structure (amorphous/crystalline; crystallographic phase), surface morphology, electron density, HfO 2 /Si interfacial layer formation, and impurity content [7-10]. Figure 1 shows the grazing incidence X-ray diffraction (XRD) spectra of 10 nm HfO 2 films grown at various temperatures using the HfCl 4 and H 2 O precursor combination. Films grown at temperature 300 C are mostly amorphous, while at 350 and 375 C are crystalline in a mixture of the orthorhombic and monoclinic phases. Surface roughness σ (measured by atomic force microscopy) and film electronic density ρ increase with growth temperature, reaching a value of σ=0.8 nm and ρ= 2.45±0.05 e - /Å 3 for film grown at 375 C. Moreover, 56

58 Figure 2. EOT values measured for HfO 2 films grown at 375 C using various precursor combinations. Figure 1. Grazing incidence XRD spectra of 10 nm HfO 2 films grown at various temperatures using the HfCl 4 and H 2 O precursors. increasing the growth temperature from 150 C to 375 C allows to reduce impurity content [10], to increase the dielectric constant from 12 to 17-18, and reduce the interface trap density (D it ) of more than one order to magnitude down to (5-8)x10 10 ev -1 cm -2. Varying the precursor combination, at a fixed growth temperature of 375 C, allows tailoring the structural/ chemical properties, in particular the surface roughness, film crystallinity, impurity content (especially carbon) and interface trap density. The film electron density varies in the e - /Å 3 range, i.e. within the range of values reported in the literature for HfO 2 dielectric. No significant variation of the HfO 2 dielectric constant and EOT of the HfO 2 /Si interfacial layer are observed for films grown using various precursor combinations (Figure 2), except for films grown using as oxygen source. For the latter, the increase of the film EOT value can be ascribed to the formation of a thick interfacial layer (SiO 2 ), as measured by cross-sectional TEM. The influence of the oxygen source (H 2 O or ) on the HfO 2 /Si interface properties was further analysed using electrically detected magnetic resonance (EDMR). This technique allows identifying the microstructure of the electrically active interface defects by monitoring spin dependent changes in the photocurrent [11]. Figure 3 shows the H II [011] EDMR spectra of two 10 nm thick HfO 2 films grown on H-terminated Si at 375 C using the HfCl 4 +H 2 O (top) and HfCl 4 + (bottom) precursor combination. HfO 2 films grown using H 2 O exhibits a Figure 3. EDMR spectra of 10 nm thick HfO 2 films grown on HF-last Si using the HfCl 4 +H 2 O (top) and HfCl 4 + (bottom) precursor combinations. Si/SiO 2 like interface, while the use of introduces a modification in the defect symmetry, changing the main axes of the interface defects wave-function. Capacitance and conductance analyses evidenced also that the interface traps density (D it ) for film grown using is one order of magnitude higher than the one measured for films grown using H 2 O. The evolution with annealing of the defect microstructure and D it was also investigated. Post growth annealing in N 2 at T>500 C allows to recover the Si/SiO 2 -like interface in terms of defect microstructure and D it. Finally, we investigated the transient currents in Al/HfO 2 /Si capacitors, considering their dependence on electric field, temperature and gate stack composition [12,13]. The HfO 2 films for this experiment were 5-57

59 Current density [A/cm 2 ] V G ON OFF I G t Open: 300K Filled: 75K x 4 J s [A/cm 2 ] Observed Region time time Time [s] Figure 4. (Top) Schematic diagram illustrating the method used for transient current measurement. (Bottom) Discharge transients measured at 300 and 75 K after positive V G charge (t on : 1000 s). In the inset the temperature dependence of stationary current at V G = 1 V. 10 nm thick and grown, using HfCl 4 and H 2 O, on chemically oxidized Si (the measured EOT of the HfO 2 / Si interfacial layer after growth is 0.9 nm) or on 2 nm thick SiO 2 layer. Reference Al/SiO 2 /Si capacitors (SiO 2 with the same HfO 2 EOT) were also measured. Transient currents were investigated according to the following procedure (Figure 4, top): a charge pulse (with amplitude V G and duration t ON ) is applied to the gate of the capacitor, then V G is set to zero and the gate current J d (discharge current) is measured as a function of time [12]. Transient currents in HfO 2 films are significantly larger than in SiO 2 (factor of 8). They slowly decrease with time following a t power-law dependence (Figure 4, bottom), and change with the amplitude of the gate pulse during the charge phase. The power-law time dependence is maintained down to the μs timescale [13] and down to 75 K. Moreover, at 75 K, J d is reduced by only a factor of 4 with respect to the values measured at 300 K. On the other hand, the static gate leakage strongly decreases with temperature (factor of 40 from 300 K to 75 K), also showing a change in the conduction mechanisms (inset of Figure 4, bottom). These results clearly t E A =87meV E A =5meV /k B T [ev -1 ] V G =0.25V V G =0.75V V G =1.25V demonstrate that transient currents in HfO 2 dielectrics do not depend on the steady-state conduction mechanisms. Dielectric polarization relaxation (DPR) [14] and charge (de)trapping (CDT) [6] are the possible origins of these slowly decreasing currents. The conduction mechanism in HfO 2 films at temperatures lower than 135 K is modelled as direct tunnelling, the extracted effective mass for HfO 2 is m=0.08 m 0 and the HfO 2 /Si barrier height is Φ=2.3eV. At room temperature trap assisted tunneling and/or Poole-Frenkel mechanism are responsible for the leakage current. In summary, the structural and electrical properties of HfO 2 films were analysed as a function of precursor combination, growth temperature and postdeposition annealing. The interface microstructure was investigated using EDMR. Transient currents in HfO 2 are significantly larger than in SiO 2 and remain an issue even at very low temperatures. Transient currents might impact on the reliability of precision circuit and memory applications. Reducing their magnitude is therefore a fundamental step for the integration of high-κ materials in advanced ICs [15]. This work was partially supported by the FIRB project n RBNE012N3X Miniaturized systems for electronics and photonics ( ), funded by Italian Ministry of University and Research. [1] G. D. Wilk, R. M. Wallace and J. M. Anthony, J. Appl. Phys. 89, 5243 (2001). [2] J. Robertson, Solid-State Electron. 49, 283 (2005). [3] V. V. Afanas ev, A. Stesmans, F. Chen, X. Shi, S. A. Campbell, Appl. Phys. Lett. 81, 1053 (2002). [4] L. Kang, B. H. Lee, W.-J. Qi, Y. Jeon, R. Nieh, S. Gopalan, K. Onishi, J. C. Lee, IEEE Electron Device Lett. 21, 181 (2000). [5] Y.-S. Lin, R. Puthenkovilakam, J. P. Chang, Appl. Phys. Lett. 81, 2041 (2002). [6] Z. Xu, L. Pantisano, A. Kerber, R. Degraeve, E. Cartier, S. De Gendt, M. Heyns, and G. Groseneken, IEEE Trans. Electron Devices 51, 402 (2004). [7] G. Scarel, S. Spiga, C. Wiemer, G. Tallarida, S. Ferrari, and M. Fanciulli, Mat. Sci. and Eng. B 109, 11 (2004). [8] M. Fanciulli, S. Spiga, G. Scarel, C. Wiemer, G. Seguini, G. Tallarida, Mat. Res. Mat. Res. Soc. Symp. Vol 786, E6.14 (2004). [9] C. Wiemer, S. Ferrari, M. Fanciulli, G. Pavia, L. Lutterotti, Thin Sol. Films., 450/1, (2004). [10]G. Scarel, C. Wiemer, S. Ferrari, G. Tallarida, and M. Fanciulli, Proc. Estonian Acad. Sci. Phys. Math. 52, 2, [11]S. Baldovino, S. Spiga, G. Scarel, and M. Fanciulli, submitted to APL. [12]C. M. Compagnoni, A. S. Spinelli, A. Bianchini, A. L. Lacaita, S. Spiga, G. Scarel, C. Wiemer, M. Fanciulli, Appl. Phys. Lett 89, (2006). 58

60 [13] C. M. Compagnoni, A. S. Spinelli, A. Bianchini, A. L. Lacaita, S. Spiga, M. Fanciulli, Microelectronic Engeneering 83, 1927 (2006). [14] J. Jameson, P. Griffin, A. Agah, J. Plummer, H.-S. Kim, D. Taylor, P. McIntyre, and W. Harrison, in IEDM Tech. Dig., pp (2003). [15] C. M. Compagnoni, A. S. Spinelli, A. Bianchini, A. L. Lacaita, S. Spiga, and M. Fanciulli, Proceedings of IRPS 2006, pp (2006). 1.4 Nanoscale electrical properties of HfO 2 and ZrO 2 thin films studied by conducting atomic-force microscopy images are used to evaluate the electrical homogeneity of the films and how it compares to their structural properties. The HfO 2 films were deposited on (001) n-type silicon wafers by ALD, using alternating pulses of HfCl 4 and water, at 300 C substrate temperature. The native oxide was not removed before growth. The ZrO 2 films were grown in similar conditions on (001) p-type silicon wafers, using ZrCl 4 and water. For both oxides, different film thicknesses, in the 5 20 nm range, were obtained by varying the number of ALD cycles only. Table I summarizes the thickness, measured by crosssectional TEM analyses, and the surface and interface roughness values, derived from AFM analysis and XRR spectra, respectively[3]. S. Kremmer 1, H. Wurmbauer 1, C. Teichert 1, G. Tallarida 2, S. Spiga 2, C. Wiemer 2, and M. Fanciulli 2 1 Institute of Physics, University of Leoben, Austria 12 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Milan, Italy Due to the wide range of possible film structure and composition, depending on film thickness and annealing conditions, the investigation of nanoscale properties in high-κ materials is of great technological importance. However, nanoscale electrical properties cannot be accessed by conventional current voltage (I V) and capacitance voltage characterization. With this respect, an already well established technique to study electrical properties on a lateral scale of only a few nanometers is conducting atomic-force microscopy (C-AFM) [1, 2 and references therein]. In this work [2], C-AFM is applied to study the electrical properties of HfO 2 and ZrO 2 thin films of various thicknesses. The local I-V curves and the 2D current ZrO 2 HfO 2 Film Surface Interface thickness (nm) roughness (nm) roughness (nm) 4.2± ± ± ± ± ± ± ± ± ± ± ± ± ± ± ± ± ±0.1 Table I. Film thickness, surface and interface roughness of measured samples. Figure 1. Grazing incidence X-ray diffraction spectra for the ZrO 2 and HfO 2 films. Diffraction peaks and relative intensities of monoclinic and tetragonal phases of ZrO 2, as well as monoclinic and orthorhombic phases of HfO 2 are also shown. From grazing-incidence XRD spectra, reported in Figure 1, the thinner films result amorphous. However, few disperse crystallites in the amorphous matrix are observed in plan view TEM images [2], which are not enough to produce a signal in the XRD spectrum. At higher thickness, films become polycrystalline with a combination of monoclinic and tetragonal phases for ZrO 2, and of monoclinic and orthorhombic phases for HfO 2. Actually, the crystalline fraction in HfO 2 films becomes large enough to be detected by XRD only in the 20 nm thick film; at the intermediate thickness only TEM analysis reveals the presence of few crystallites dispersed in the amorphous matrix. The surface roughness in both materials increases with film 59

61 thickness and crystalline fraction: in agreement with the behavior of the structural properties, the increase is almost linear for ZrO 2, whereas in HfO 2 it becomes significant only in the 20 nm film [2]. The C-AFM experiments were performed under UHV conditions, in contact mode using boron-doped diamond-coated silicon probes. For detailed studies on the electrical characteristics, local I V curves are recorded as well as two-dimensional (2D) tunneling current images. For the local I V characterization, the current through the tip is measured during the application of negative voltage ramps to the sample. From these I V curves a pseudo-breakdown voltage (PBDV) is extracted, which is arbitrarily defined as the voltage where a current of 1 pa is reached in the measurement. The 2D tunneling current images are obtained by applying a constant positive voltage to the sample and recording the current through the tip, whilst scanning for the topography. A comparison of the PBDV characteristics for the different oxide films is summarized in Figure 2. The number of counts where the current in an I V curve exceeded 1 pa (PBDV) is plotted versus the voltage reached at this point. The total number of measurements for each oxide is indicated in the upper right corner of the individual graphs. The mean applied voltage of the PBDV distribution, for similar oxide thickness, is always higher in the case of the HfO 2 films. This is due to the different doping type of the substrate Figure 2. Number of PBDV counts vs applied sample voltage for the ZrO 2 and HfO 2 films. The numbers in the upper right corner of each graph indicate the total number of measured I-V curves. Figure 3. Topography (a) and 2D current image (b) measured on the 8.2 nm ZrO 2 film. The imaging voltage was 5.7 V. The z scale is 4 nm in the topography image and 5 pa in the current image. Topography (c) and 2D current image (d) measured on the 10.5 nm HfO 2 film. The imaging voltage was 8.8 V. The z scale is 2.5 nm in the topography image and 5 pa in the current image. The circles and the corresponding numbers mark areas with distinctive current and topographical characteristics. that shifts the tunneling current onset to more negative voltages in the case of the HfO 2, as well as to its slightly higher thickness. The PBDV distribution for the 4.2 nm ZrO 2 is rather narrow, with a small side peak visible at higher voltages. This side peak could be explained by the presence of crystallites in the 4.2 nm ZrO 2 film as they can be observed in the plan view TEM images. The 8.2 nm ZrO 2 film exhibits a slightly broadened PBDV distribution compared to the 4.2 nm film, with the side peak now merging with the main distribution. This behavior can be attributed both to the increased surface roughness and to an increase in the crystalline fraction, as determined from XRD. In fact, assuming that in both films these crystallites are in the tetragonal metastable phase, whose calculated dielectric constant is very high [4], a higher voltage is required to measure 1 pa from these sites. The statistical PBDV distributions of the 5.1- and the 10.5-nm HfO 2 samples are very similar in width and do not present any side peak, in agreement with the mostly amorphous nature of these films. However, a closer look at the 10.5-nm HfO 2 graph reveals some counts which deviate from the narrow distribution. It is feasible that these outliers are related to the very small fraction of crystallites detected by the TEM analysis, as well as to the contribution of grain boundaries, which are expected to be more conductive than the 60

62 matrix. For the two thickest films the obtained PBDV distributions are strongly broadened. One reason for this is the increase in the surface roughness, although it cannot be the only parameter involved. In fact, the ZrO 2 film exhibits a higher rms roughness compared to the HfO 2 film. In contrast to this, the distribution is much more broadened in the HfO 2 case. Since both films are mainly polycrystalline, different crystallite structures and orientations, together with the presence of grain boundaries contribute to the broadening of the PBDV distribution. In Figure 3, the 2D current images for the intermediate thickness of the two oxide films are presented, along with their morphology image. Comparing the 2D current images, it results that the current spots in HfO 2 are larger. However, their number is higher in ZrO 2. This is valid also at the other thickness values [2]. We preliminarily observe that the side peak present in the PBDV distribution for ZrO 2, and assigned to the presence of crystallites in the film, is not expected to play much role in the image reported in figure 3, since the latter is recorded at a lower voltage. locations the oxide leakage can be attributed to thickness inhomogeneity due to interface roughness. Only these two types of leakage spots are observed in the ZrO 2 films, therefore it can be stated that the electrical 2D characteristics of the 4.2 and the 8.2 nm ZrO 2 films are mainly determined by thickness inhomogeneity. In the HfO 2 film, four different locations marked by numbers can be distinguished. The spots marked by 3 show a peak in the topography and in the current image. The topographical peak could be a crystallite growing higher into vertical direction and having higher conductance than the amorphous matrix. This behavior could be attributed to the fact that some crystallites might have a lower-κ value compared to the amorphous matrix. Label 4 marks a location where the topography shows a protrusion but no measurable current is visible in the current image. This protrusion could also be due to the growth of a crystallite, but with lower conductance compared to the one in spot 3. Label 5 marks a deep depression in the topography image. As its depth of about 1 nm is much larger than the surface and the interface rms roughness, a significant thinning of the sample can be assumed at this spot. The fact that still no current is measured indicates that the influence of crystallite properties in 3 is much more severe than the influence of thickness inhomogeneities (1/2/5). Label 6 marks a high current spot, which is close to a protrusion. If this protrusion is due to a crystallite, the increased current could be caused by a grain boundary at this location. In general, it can be stated that the weak spots in the current images of the 5.1 and the 10.5 nm HfO 2 are mainly caused by the presence of crystallites in the films. This work is partially supported by the Ministry for Foreign Affairs, in the framework of the scientific and technological collaboration between Italy and Austria, project 9/2004. Figure 4. Schematic illustration of the different phenomena influencing the leakage current distribution observed in figure 3. Leakage current spots can be related to several structural and morphological characteristics. In Figure 4, some of these are sketched and related to features of Figure 3. The locations marked with 1 indicate depressions in the topography correlated to high currents in the current image. Here, a thickness inhomogeneity due to the surface roughness of the amorphous film seems to be the major contribution to the electrical weakness. Label 2 indicate high current spots, where no depression in the morphology is observed. In these [1] M. P. Murell, M. E. Welland, S. J. O Shea, T. M. H. Wong, J. R. Barnes, A. W. McKinnon, M. Heyns, and S. Verhaverbeke, Appl. Phys. Lett. 62, 786 (1993). [2] S. Kremmer, H. Wurmbauer, C. Teichert, G. Tallarida, S. Spiga, C. Wiemer, and M. Fanciulli, J. Appl. Phys., (2005). [3] C. Wiemer, S. Ferrari, M. Fanciulli, G. Pavia, and L. Lutterotti, Thin Solid Films 450, 134 (2004); S. Ferrari, D. T. Dekadjevi, S. Spiga, G. Tallarida, C. Wiemer, and M. Fanciulli, J. Non-Cryst. Solids 303, 29 (2002). [4] X. Zhao and D. Vanderbilt, Phys. Rev. B 65, (2002). 61

63 1.5 Thermal stability of HfO 2 /TiN gate stacks for 45 nm CMOS devices C. Wiemer 1, M. Perego 1, M. Fanciulli 1, V. Cosnier 2, P. Besson 2, V. Loup 3, L. Vandroux 3, S. Minoret 3, M. Cassé 3, X. Garros 3, J-M. Pedini 2, S. Lhostis 2, K. Dabertrand 2, C. Morin CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy STMicroelectronics, Crolles Cedex, France CEA/LETI, Grenoble Cedex, France The implementation of high-κ dielectrics in logic devices requires the introduction of metal gates with a suitable work function. For the mostly studied high-κ candidate, HfO 2, TiN metal gates are proposed due to their high work function. Here we evaluate the impact of the gate stack layers deposition technologies and their combination on the thermal stability of the stack with respect to the equivalent oxide thickness (EOT) figure of merit. Two HfO 2 deposition technologies have been used: ALCVD (atomic layer chemical vapor deposition) and AVD (Atomic Vapor Deposition). ALCVD process is carried out from HfCl 4 and H 2 O precursors at 350 C, with a post deposition annealing at 600 C, whereas the AVD process involves Hf(OtBu) 2 (mmp) 2 and oxygen at 550 C, with no post-deposition annealing. The nominal thickness was 2 nm in both cases. For this thickness value, both deposition techniques are known to provide layers that remain amorphous even after annealing at high temperature [1]. Two technologies for the deposition of 10 nm TiN on the gate dielectric have been evaluated: CVD, from TiCl 4 and NH 3 at 620 C, and PVD at 300 C. The full thermal budget of a complete process flow was simulated by a spike anneal at 1050 C after the deposition of poly-silicon on top of the metal gate. The structural and morphological evolution of the gate stacks was studied by X-ray reflectivity (XRR). This technique evidenced the much better stability of PVD TiN/AVD HfO 2 stack as compared to the other stacks. This is illustrated in Figure 1, where XRR data, simulations and corresponding electronic density profiles are reported. Indeed, when PVD TiN is used, interfaces are much sharper and there is much less variation after anneal. In particular, the TiN/Si interface is much rougher for CVD than for PVD TiN (1.7 vs 0.7 nm), and the difference is Figure 1. (a) polysi/cvd-tin/alcvd-hfo 2 before (blue) and after (red) 1050 C spike anneal. The same for: (b) polysi/pvd TiN/HfO 2 ALCVD, (c) polysi/cvd TiN/HfO 2 AVD, (d) polysi/pvd TiN/HfO 2 AVD. 62

64 in the ALD HfO 2 layer (Figure 2(a)). No chlorine has been observed in the PVD TiN films and in the AVD HfO 2 layers (Figure 2(c)). The CVD TiN films are characterized by a slightly higher oxygen content compared to the PVD ones (data not shown). The oxygen profile remains unchanged after 1050 C spike anneal. The analysis of the HfN - signals after thermal treatment indicates Hf diffusion in the TiN film. Higher HfN - signals are observed in the CVD TiN films compared to the PVD ones (Figure 2(b) and 2(d) respectively). This result correlates with the higher roughness at the TiN/Si and TiN/HfO 2 interfaces and to the higher electronic density variation evidenced by XRR. Indeed, since the electronic density of HfN x or HfO x N y is higher than the one of TiN, inter-diffusion of these species into the TiN layer can explain the observed increase in the electronic density. The diffusion process seems not to be influenced by the characteristics of the underlying dielectric. Figure 2.ToF-SIMS depth profiles of the CVD TiN layer deposited on the ALCVD HfO 2 film before (a) and after annealing (b). Similar data for the PVD TiN layer deposited on the AVD HfO 2 film before (c) and after annealing (d). Figure 3 shows the measured EOT for these metal gate/ high-κ dielectric combinations. The variations in EOT correlate with the variation of interfacial roughness as measured by XRR. even higher after the spike anneal (2.3 vs 0.7 nm). The HfO 2 /TiN interface is flatter than the TiN/Si, although annealing induces a roughness increase, especially in the case of HfO 2 grown by ALCVD. It is also interesting to note how the thickness of the HfO 2 layer considerably reduces after spike anneal when CVD TiN is used. This roughening can be related to both physical roughening and to some chemical intermixing at the interfaces. The electronic density of CVD TiN is found to increase after annealing. Grazing incidence X-ray diffraction (GIXRD) measurements show only small differences between PVD TiN and CVD TiN, with a broader peak at 2θ ~ 42 for PVD TiN. This is indicating a possible lower crystallization state or higher microstrain, probably related to the lower deposition temperature of PVD than of CVD TiN. No effect of the underlying dielectric has been observed on the TiN microstructure. After 1050 C spike anneal, the silicon cap layer crystallizes and all diffraction pattern are the same (not shown). Therefore, the variation of the electronic density of CVD TiN cannot be related to modification of its crystallographic ordering. Moreover, since variations of the electronic density are revealed in the HfO 2 layer as well, interdiffusion phenomena are very probable. Tof-SIMS analysis indicates the presence of very high chlorine contaminations in the CVD TiN films and Figure 3. EOT measured on HfO 2 /TiN capacitors as a function of deposition technology and thermal budget. In PVD TiN/AVD HfO 2 stacks better thermal stability and lower EOT correlate with the lower contaminant content, which may lead to lower TiN/HfO 2 interfacial roughening. Combination of AVD HfO 2 + PVD TiN therefore ensures the stability of the interfaces. This work is achieved within the framework of the MEDEA+ Foremost program. [1] V. Cosnier, P. Besson, V. Loup, L. Vandroux, S. Minoret, M. Cassé, X. Garros, J-M. Pedini, S. Lhostis, K. Dabertrand, C. Morin, C. Wiemer, M. Perego, and M. Fanciulli, submitted. 63

65 1.6 Oxygen diffusion in HfO 2 /SiO 2 /Si stacks S. Ferrari and M. Fanciulli I 18 /(I 18 +I 16 ) SiO 2 Si Stage 1 Stage 2 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy Sputter time (s) The nature and the properties of the HfO 2 /Si interface have a number of implications on the electrical and physical properties of the oxide. In particular, the mechanism of atomic exchange at the interface, especially for oxygen, is of particular interest from a fundamental perspective, as well as for its technological implications related to the formation of the interfacial oxide that determines the performances of the dielectric stack in the high-κ based CMOS technology. Given the difference in the oxygen diffusion mechanism between the SiO 2 and the HfO 2, the silicon oxidation process in the presence of an HfO 2 layer can occur either through a complicated oxygen exchange mechanism at the HfO 2 /SiO 2 interface or via the diffusion of either atomic oxygen into silicon or molecular oxygen in HfO 2. The mechanism of oxygen diffusion in SiO 2 in the presence of HfO 2 has been the subject of intense investigation in our group [1]. By annealing HfO 2 /SiO 2 stacks in 18 O 2 atmosphere and by studying the 18 O profile into the SiO 2 by Time of Flight Secondary Ion Mass Spectrometry (ToF-SIMS) it was possible to elucidate the oxygen diffusion/exchange mechanisms. In general the total amount of 18 O uptake in the SiO 2 layer underneath the HfO 2 is very limited as compared to the bare SiO 2 and accumulates at the SiO 2 /Si interface as in the case of bare SiO 2. This is quite surprising, since such behavior suggests that molecular 18 O 2 takes part in the diffusion process in the HfO 2 /SiO 2 /Si stack. The mechanisms occurring during O 2 annealing can therefore be divided in two stages, the first one (stage 1) is already complete after 10 s, whereas the second one (stage 2) starts afterwards and follows a much slower kinetics. We determined the 18 O profiles in the capped SiO 2 layer of stage 1 and stage 2, by estimating the 18 O ToF-SIMS profiles differences of various samples. The results are shown in Figure 1. We can see that Figure O concentration profiles of early oxidation stage (stage 1) and late oxidation stage (stage 2). The profiles were offset placing the HfO 2 /SiO 2 at the origin. (See text for a detailed explanation.) after stage 1 oxidation, the resulting 18 O distribution is almost flat throughout the SiO 2 layer. The dip in proximity of the HfO 2 /SiO 2 interface is a measurement artifacts. The 18 O profile for the stage 2 oxidation, on the other hand, is surprisingly similar to the 18 O profile of the annealed bare SiO 2. We make the hypothesis that stage 1 oxidation is promoted by moisture adsorbed in the HfO 2 layer. Such moisture is unstable at high temperatures during annealing in oxygen, therefore its removal produces changes in the oxidation kinetics. We monitored the OH and H signals in the TOF-SIMS depth profile to get an indication of the trend of the moisture desorption kinetics as a function of annealing temperature and atmosphere. A ToF-SIMS analysis of a set of samples annealed at 500, 700 and 900 C for 10 s in O 2 and N 2 respectively, is shown in Figure 2. Both H and OH peaks provide an indication of the presence of hydrogen in the high-κ layer. We can observe how the annealing in O 2 is extremely effective in removing hydrogen. Indeed, the instrumental background level is reached at 700 C. Moisture is adsorbed preferentially at the HfO 2 /SiO 2 interface, as it can be concluded observing the peak at the interface. For samples annealed in N 2, a OH and an H peak is still present in the sample annealed at 900 C, demonstrating that moisture removal upon annealing in N 2 has little effect. We can therefore refer to stage 1 oxidation as a wet oxidation process. It is well known that HfO 2 is much more hygroscopic than SiO 2 and also that moisture adsorption is partially irreversible in HfO 2. Oxygen exchange rate into the HfO 2 is extremely fast, while on the other hand it appears that HfO 2 constitutes a formidable barrier for the oxidation of the underlying silicon. Our 18 O diffusion studies demonstrate that the species responsible for stage 2 oxidation is molecular oxygen. If the oxygen leading to stage 2 derives from the HfO 2 lattice there must be an oxygen exchange process at the HfO 2 /SiO 2 interface 64

66 Intensity (Arb. u.) HfO 2 SiO 2 (a) SiO 2 (1) (2) (3) (4) Si the oxidation process is the same for the HfO 2 /SiO 2 and the bare SiO 2 indicating that the rate determining step for the oxidation is the same in the two cases. We therefore conclude that the stage 2 oxidation is the result of molecular oxygen diffusion and that the atomic oxygen moving fast into the HfO 2 lattice cannot exchange with the SiO 2 at the interface Intensity (Arb. u.) X(nm/hr) Δ Sputter time (s) Sputter time (s) Figure 2. OH signal of O 2 (a) and N 2 (b) annealed samples. Samples as-deposited (1 [black]) and annealed at 900 C (2 [red]), 700 C (3 [green]) and 500 C (4 [blue]). OH signals gives an indirect information on moisture adsorbed in the HfO 2 film. While for O 2 annealing, 700 C are sufficient to eliminate excess moisture at the interface, in N 2 excess OH is found even in sample annealed at 900 C HfO 2 SiO 2 SiO 2 SiO 2 HfO /T (K -1 ) Figure 3. Arrhenius plot of the oxidation velocity of SiO 2 in a HfO 2 /SiO 2 (close circles) stack and on a bare SiO 2 (open squares) films as a function of temperature. transforming charged-atomic oxygen into neutralmolecular oxygen. Figure 3 shows the Arrhenius plot of the oxidation velocity for the two considered systems. We can see that the slope of the two lines is the same, suggesting that the activation energy for (b) (1) (2) (3) (4) Si [1] S. Ferrari and M. Fanciulli, Journal of physical chemistry B, 110, (2006). 1.7 Characterization of the mechanical stress induced in silicon during device fabrication E. Bonera 1, M. Fanciulli 1, G. Carnevale 2, M. Mariani 2 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy 12 STMicroelectronics, Agrate Brianza, Italy Raman spectroscopy can be effectively used for the characterisation of strain in silicon crystals with a diffraction limited spatial resolution and a sensitivity to stress lower than 100 MPa. Nevertheless, the power of Raman spectroscopy (RS) as a tool for local mechanical stress determination in microelectronics was strongly reduced when the industrial design rule became smaller than the diffraction barrier of half the excitation wavelength λ. The first problem arisen has been of course the spatial resolution, but, although limiting, this factor is not fundamental. Indeed, in many cases from the applicative point of view it is not necessary to spatially resolve individual devices, but rather to obtain information of the stress state in large arrays of structures. The truly limiting factor of RS for subwavelength structures is averaging, as collecting a Raman spectrum from an inhomogeneously strained volume greatly reduces the sensitivity of the technique. Another problem of averaging is also the limitation to the interpretation of the experimental results. The need for the unparalleled advantages of RS (non-destructivity, speed, reliability, clean-room 65

67 compatibility, cost) pushed the research towards the idea of employing an ultra-violet excitation to transform a volume analysis into a surface analysis, thanks to the strongly reduced penetration depth d P. The higher sensitivity is a consequence of the less limiting impact of a bidimensional averaging with respect to volume averaging. a nearly year basis. In addition, in the last years there has been an effort to intentionally introduce strain in order to enhance the performances of devices. Recent experimental results, obtained also on nonuniform subwavelength devices, represent a strong motivation in using this spectroscopy to validate the FE modeling techniques extensively employed in microelectronics. With the results presented in Figure 1, we showed that, although diffraction limited, Raman spectroscopy can be an effective tool for the measurement of strain in subwavelength-patterned silicon devices, if the advantages of a resonant excitation and the transmission properties of a grating are simultaneously exploited [1,2]. Analyzing different steps of an industrial process we could monitor qualitatively, but effectively, the evolution of the strain. An example of stress evolution is reported in Figure 2. We were able to single out the crucial process step for the formation of defects that could affect the performances of the final device. As the technique takes advantage from the shrinking of the measured structures, we believe it will be also a valuable tool for the characterization of the future microelectronic technology. Another drawback of Raman spectroscopy is the difficulty of obtaining quantitative information on the stress tensor from the shift of the phonon bands. Whereas the technique can in principle access all the information about the stress σ by simply setting the appropriate experimental geometries, the practical interpretation of μrs maps is always complicated mainly by two factors. One is that the tensorial nature of σ can be practically exploited by μrs only in peculiar cases, while in most experiments it is necessary to make assumptions on the shape of the tensor. The other is that there is an averaging effect when the structures investigated are inhomogeneously stressed on a spatial scale of the order of the excitation wavelength λ. In many cases it is therefore impossible to univocally correlate to σ the experimentally determined value of Δω exp, i.e. the variation of phonon energy in the strained silicon crystal. For this reason some authors adopted an interpretation approach based on the comparison of the experimental data with the strain calculated with a finite-element (FE) model. We developed a virtual experiment [3,4] that reproduces completely the whole process of data acquisition, in which also many of the instrumental parameters and experimental conditions are taken into account. The procedure was adopted to study the case of stress in microelectronics devices. This topic was also of practical interest because the everdecreasing dimensions of devices revive stress-related problems at every change of the technological node on Figure 1. Raman maps of five different sets of test structures, labeled from A to E. The scheme of the alternation of active areas (dark) and isolations is drawn under the maps. (inset) Sketch of the cross-section of shallow trenches and active areas. Figure 2. (a) Evolution of the stress with the advance of the manufacturing process for the test structures E reported in Fig. 1. (b) Evolution of the stress of 100 nm memory devices. In panel (b), the data for steps from 6 to 10 are missing as the structures are masked by an opaque layer. 66

68 These works determined also the choice to focus on the ultraviolet (uv) μrs. It has been shown that the experimental conditions of this technique are incompatible with some of the assumptions made in the previous literature. The method we employed is based on the idea of simulating an experiment in which we take into account also many details of the instrument and the data acquisition geometry. Figure 3 gives an idea of the simulation procedure. any parameter on the base of the experimental results themselves. The simulation could of course be improved with threedimensional FE modeling and rigorous electromagnetic field calculations. Nevertheless, when applied to a microelectronic device manufacturing problem, this method validated the finite-element modeling in most of the cases considered. In addition, we showed that the process leading from non-uniformly strained structures and a spreaded distribution of Δω α,j to a single value Δω exp is particularly complicate. This difficulty is reflected in the quantitative analysis of data when the FE simulation is not available. [1] E. Bonera, M. Fanciulli, M. Mariani, Raman spectroscopy of strain in subwavelength microelectronic devices, Appl. Phys. Lett. 87, (2005). [2] E. Bonera, M. Fanciulli, Resonant Raman microscopy of stress in Silicon-based microelectronics, Microscopy of Semiconducting Materials XIV (2005). [3] E. Bonera, M. Fanciulli, G. P. Carnevale, Raman stress maps from finite-element models of silicon structures, J. Appl. Phys. 100, (2006). [4] P. Fantini, A. Ghetti, C.P. Carnevale, E. Bonera, D. Rideau, A full self-consistent methodology for strain-induced effects characterization in silicon devices, IEDM Tech. Dig. 2005, 992 (2005). Figure 3. Upper Panels: (a) A scanning electron micrograph of the structure and (b) the finite-element grid used to model it. Lower Panels: (a) Illumination of the structure when the focus is positioned at an external edge of 2.0 μm structures. The lateral extension of the focus is wider than the magnification of this panel. This illumination produces the spectrum reported in panel (b). The symbol Σ indicates that this spectrum is the result of the summation of all the finite-element contributions. Panels (c) and (d) show two single finite-element spectra (solid) generated by elements located at a convex and a concave angle, respectively. The two components (dotted) show the splitting induced by shear strain. The central vertical line marks the unstrained silicon value. In particular, the most important innovations introduced in this algorithm are an illumination which is dependent on the shape of the silicon structures, the simulation of illumination and collection with a finite angle of observation, and the contribution of all the Δω expα,j to the final spectrum. These features are of particular importance when the experimental maps are acquired with a small d P like in the case of a silicon-resonant excitation, where the reduced averaging increases the sensitivity but also the inhomogeneity and complexity. In order to use this method as a proof of the goodness of the FE modeling, we developed it without adjusting 1.8 Low-κ materials for intra-metal dielectrics A. M. Ferretti, C. Wiemer, E. Bonera, C. Rossi, M. Fanciulli CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy In order to reduce the cross-talk between interconnects, lower κ dielectrics are needed. Recently the attention has focused on inorganic-organic films like silicon oxide incorporating methyl groups (SiOCH) because they can assure, together with a low dielectric constant, good thermal stability and adhesion to other layers. However, the chemical structure of SiOCH films with respect to the deposition and processing conditions is still unknown. Moreover, 67

69 the determination of the film porosity becomes a fundamental issue because this parameter strongly influences the dielectric constant. Using three different techniques, X-ray scattering, Raman and EPR spectroscopy, we obtained information on the chemical structure of these films that allowed us to make some hypotheses on the influence of the stochiometry on the film porosity. We measured three different films STD (standard preparation), HP (high pressure process), and PETEOS, which is the SiO 2 reference sample, deposited at 350 C by Plasma Enhancement Chemical Vapor Deposition (PE-CVD) by STMicroelectronics. The characteristics of the films are reported in Table I. Refraction Dielectric index (633nm) constant κ STD HP PETEOS Table I. Characteristics of the studied film. The Raman spectrum of the STD sample shows only a band at 1446 cm -1, corresponding to the graphite C-C bond while the HP one shows the same band split into two features centered at 1412 and 1446 cm -1 indicating a different carbon surrounding, and also two bands in the spectral range from 2900 to 3000 cm -1 corresponding to C-H bond frequencies. We studied these two samples also using EPR spectroscopy at 34 GHz. This high frequency allows resolving contributions due to different species having close g factor. We have conducted a systematic study by recording spectra after a drying treatment at 100 C in N 2 flow for 15 minutes, and after thermal treatments which reproduce the one done during the process, where the maximum temperature achieved is 400 C. The samples have been also studied before and after irradiation with X-ray at 30 kv and 32 ma for 4 hours. Two different signals have been resolved, one at g = ± and the other one at g = ± : the first is due to the silicon dangling bond (Si-db) in amorphous matrix and the second one is due to the carbon dangling bond (C-db) [1]. The different nature of the two signals is confirmed by the analysis done as a function of the microwave power (Figure 1). Indeed, the intensity of the lower field signal increases with the microwave power, while the higher field one is less sensitive to the microwave power. The C-db signal is normally more Figure 1. EPR spectra of HP as deposited after X-ray irradiation film recorded at room temperature at 34 GHz, as a function of the microwave power. intense in the irradiated sample, because the X-rays easily break the carbon bonds leaving a radical defect in the film. Although these signals have not any hyperfine interactions, which could have been useful to characterize the chemical environment of the defect, they provide interesting information from a qualitatively point of view on the film structure. In fact, in the HP the Si-db is not detected and the C-db is really strong, while in the STD sample both signals are present. On the basis of these observations it is possible to assert that in the HP sample the number of broken bonds is higher, suggesting a higher number of alkyil groups. The electronic density of these materials is evaluated from the critical angle of specular X-ray reflectivity (XRR) data. The electronic density of STD SiOCH is in the e - /Å 3 range, whereas a value of 0.46 e - /Å 3 is obtained with the HP process. Detector scans at fixed low incidence angle have been performed. Due to the low surface and interfacial roughness (~6 Å and ~2 Å, respectively), the diffuse scattering from pores dominates over that from roughness for angles away from the critical angle. The incidence angle was then chosen between the critical angle of the low-κ and the critical angle of the silicon substrate, so that the radiation is not penetrating into the silicon substrate. In order to extrapolate the porosity and the pore diameter from the diffuse scattered X-ray intensity, we simulate the X-ray intensity using the distorted-wave Born approximation (DWBA). Fixing the electron density of the skeleton of the low-κ, the porosity (defined as P= 1-SiOCH density/skeleton density) can be determined by the critical angle of the specular XRR curves. The electron density of the skeleton is assumed to be equal to either amorphous SiO 2 (2.20 g/cm 3 ), or Hydrogen Sesquoxiane (HSQ) 68

70 (2.00 g/cm 3 ), or Si0 1.5 C 1.2 H 3, giving: 1.59 g/cm 3 for STD and 1.46 g/cm 3 for HP. Knowing the mass density of STD SiOCH to be equal to 1.35 g/cm 3, and the electronic density of STD and HP SiOCH as measured by XRR, and assuming that the difference of the electronic density is related to variation of porosity and not to variations of the skeleton, the mass density of HP SiOCH is equal to 1.24 g/cm 3. The porosity and pore size as obtained assuming a random two-phase model in which one phase consists of the pores and the other of the surrounding matrix material [2] are reported in Table II. the diffuse scattering curves could also be related to a different chemical composition influecing the calculation of the porosity value. HP SiOCH is then less dense, with higher pore size and higher density of methyl groups than STD one. The lower κ value of HP SiOCH corresponds then to the less dense film with a higher content of methyl groups. In view of an integration of these low-κ materials, we also investigate their evolution when exposed to various plasma treatments necessary for the photoresist removal process step. For 40 nm thick films, O 2 treatment causes a thickness STD SiOCH HP SiOCH SiO 2 HSQ SiO 1.5 C 1.2 H 3 SiO 2 HSQ SiO 1.5 C 1.2 H 3 Porosity (%) Pore Size (Å) Table II. Porosity (%) and pore size (Å) as obtained for STD and HP SiOCH using different skeletons. The pore size of HP SiOCH results to be higher than the one of STD SiOCH. However, due to the variation in the chemical structure between HP and STD films evidenced by Raman and EPR analysis, variation in reduction, together with an increase of the electronic density, as shown in Figure 2(a). The value of the electronic density ( e - /Å 3 ) is now close to the one of amorphous SiO 2 (0.67 e - / Å 3 ). The reduction in thickness is higher for HP than for STD SiOCH (32% vs 27%). For the H 2 treated samples, the electronic density, roughness, and thickness are the same, within the experimental error, of those found on as grown samples (Figure 2(b)). No modification of the SiOCH films are found after H 2 treatment by XRR analysis. In the EPR spectra the O 2 plasma treatment causes a relative decreasing of the Si-db signal with respect to the C-db (Figure 3) signal probably because of Figure 2. XRR spectra (dots) and fits (lines) of the as grown (black), O 2 (red, a), and H 2 (blue, b) treated STD SiOCH. Figure 3. EPR spectra of STD film: as deposited (black), O 2 (red), H 2 (blu), recorded after thermal treatment reproducing the process and after X-ray irradiation. 69

71 the oxidation of silicon, while, after the H 2 plasma treatment, there is no significant modification in the spectra of both samples with respect to the spectra of the as deposited film. Concluding we can assert that the HP sample is less dense and with larger pores than the STD one. Moreover, comparing the effects of the two different plasma etchings, we discovered that the O 2 one oxidizes the film, while the H 2 one leaves the film structure unchanged. [1] T. Ehara, K. Notake, K. Handa, Diamond and Related Materials, 10, (2001). [2] C.H. Russell, Mat. Res. Soc. Symp. Proc. 716, B1.3.1 (2002). 1.9 Advanced materials for interconnects G. Tallarida 1, C. Wiemer 1, L. Aina 2, S. Alberici 2, E. Ravizza 2, A. Giussani 2, G. Pavia 2, E. Varesi 2, G. Brunoldi 2, S. Guerrieri 2, S. Grasso 2, E. Ravizza 2, S. Spadoni 2 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy 2 STMicroelectronics, Agrate Brianza, Italy Minimizing RC delay has forced a transition from aluminium-copper, tungsten, and silicon dioxide based interconnects to Cu and low-dielectric constant metallization schemes. Copper is currently deposited by electroplating (ECD) techniques to fill a damascene pattern, followed by a chemicalmechanical processing (CMP) [1]. One major concern in this process is related to the high diffusivity of copper in Si and Si-based dielectrics, which requires the introduction in the interconnects of barrier stacks capable of stopping copper diffusion and of ensuring optimal adhesion between copper and the intra-metal dielectric stack. Currently PVD deposition of Ta based stacks, followed by the deposition of a copper seed layer, is the preferred route for obtaining reliable diffusion barriers and prepare for the ECD copper growth [2]. However, this approach cannot be scaled down to meet the requirements of the 45 nm technology node and beyond. In-fact, the PVD process does not allow a conformal growth, hindering its use in ultra narrow lines, where the preferential growth is likely to induce the formation of voids and/or the thinning (or even the vanishing) of the barrier at the sidewall and at the corners. On the other hand, to keep the conductivity of the line suitable for achieving the maximum signal transmission, the barrier stack material has to provide satisfactory electrical properties. To meet these requirements, several deposition techniques and materials are being investigated, focusing on those processes capable of ensuring the conformal growth (such as CVD and ALD process) of low resistivity thin films. The latter include Ta-, W-, and Ti-based binary and ternary compounds, which are the most studied owing to their desirable physical, chemical, and electrical properties. To evaluate the efficiency of the barrier stack and its thermal stability, we developed a procedure based on the study of the barrier/copper structural and chemical properties and how these evolve with thermal annealing. This is achieved by combining XRD and XRR, with ToF-SIMS and XPS analyses. In the following, the main results achieved on ALD TaN based barrier stacks are highlighted [3] and compared to recent results on WN [4]. Moreover, two aspects related to the interplay between the structural properties of ECD copper and its conductivity are addressed. Measured samples consist of ALD-TaN/PVD-Ta/ PVD-Cu stacked structures. ALD TaN was deposited by alternating pulses of metal-organic Ta precursor and NH 3, using Ar as purging gas and keeping the substrate at 275 C. The stoichiometry of the TaN layer is close to Ta 3 N 5 ; however, significant amounts of C and O are also present in the layer. Samples were cut from the same wafer with nominal 2nm TaN/5nm Ta/30nm Cu-seed stack, on which a layer of ECD-Cu, 400nm thick, was deposited and CMP finished. Samples were annealed at 200 C, 300 C, 400 C, 500 C and 600 C for 1 hour in nitrogen flux. XRR spectra of samples annealed at 200 C and 300 C fully overlap, indicating that no appreciable structural variation occurs up to 300 C. As shown by the evolution of the electronic density profile reported in Figure 1 for selected samples, significant structural changes are found after annealing at 500 C and 600 C: the thickness of the barrier layer, as well as the barrier/copper roughness, increases, whereas the electronic density decreases. The most relevant 70

72 Figure 1. Electronic density profile of TaNTa/Cu stack after annealing at various temperatures, derived from the fitting of the XRR spectra. modifications occur at the barrier/copper interface, whereas the barrier/oxide interface is not greatly affected by the annealing. This might be an indication of barrier/copper degradation, caused by interdiffusion phenomena. To evaluate whether the structural degradation is associated with copper diffusion into the oxide, the same samples were investigated by ToF-SIMS. Results are reported in Figure 2. Significant modifications of the elements distribution are already evidenced after annealing at 400 C. In fact, with reference to the Ta peak maximum, the copper signal starts to appear earlier than in the asdeposited sample, and this might indicate that the barrier efficiency towards Cu diffusion is reducing. Moreover, TaO, TaN and Ta signals have a different shape, indicating that a certain degree of chemical/ structural changing is taking place. Modifications are more pronounced after annealing at 500 C, where copper signal starts to appear almost simultaneously to barrier related signals, suggesting that interdiffusion is taking place in this sample. After annealing at 600 C the trend is much more marked. Nonetheless, the copper signal in the dielectric region close to the barrier layer is below the detection limit of the ToF-SIMS, indicating that, according to ToF-SIMS sensitivity, copper is not diffused into the dielectric layer. We conclude that ALD TaN, although very thin, is an effective barrier up to 300 C, whereas annealing at 400 C induces chemical, and to a lesser extent structural, modifications at the barrier/ copper interface. Compared to the results obtained in our previous work on PVD-TaNTa barriers [2], the investigated ALD-TaN based barrier shows a slightly weaker thermal stability. In fact, PVD-TaNTa barriers were not significantly affected by annealing up to 400 C. Differences between the two types of barrier are mostly observed in the ToF-SIMS profiles and are Figure 2. ToF-SIMS chemical profile in the copper/barrier region after various annealing temperature. Figure 3. Evolution of TaN and WN surface roughness and thickness with thermal annealing. The data are derived from the fitting of XRR spectra. related to the copper/barrier interface. However, the thickness might play an important role in defining the efficiency of the barrier and, since the investigated PVD TaNTa stacks were 12.5 nm thick, no definitive word can be said in this sense. WN layers, in a similar thickness range and produced on 200mm wafer size tool by PNL (Pulse Nucleation Layer a process similar to ALD) at 300 C using B 2 H 6, WF 6 and NH 3, were also investigated. 71

73 Compared to ALD TaN, the WN/Cu system presents a different evolution with annealing. Barrier thickness increase and electron density decrease are not as severe as in the TaNTa stack, whereas a significant increase of the barrier/copper roughness is observed, as shown in Figure 3. From ToF-SIMS analysis, it results that after annealing at 400 C the barrier region slightly broadens and the copper/barrier interface significantly roughens, whereas the barrier/oxide interface does not seem to be affected by the thermal treatment. These effects are more pronounced by the annealing at 500 C and become striking after 600 C, when the barrier results almost completely dissolved into the copper. From XRD measurements, it was observed that WN crystallizes mainly as WN after 400 C treatment, and as W 2 N after 500 C. This might be an indication of nitrogen diffusing out from WN, although ToF-SIMS profiles could not confirm this hypothesis. The storage in air of the WN layers before copper stack deposition caused the formation of an oxidized, low density, top layer. This oxidized layer is probably responsible for the bad adhesion between copper and WN observed in some samples. Bad adhesion, which was not observed in the ALD- TaN based barrier, is a major drawback for the use of WN. On the other hand, the use of either ALD-TaN or WN, will probably require a decrease of the total thermal budget experienced by the metallization, since in both cases significant structural and chemical changes are already observed in the temperature range used in the current back-end processes. Figure 4. XRD texture analysis of a copper film deposited on α-ta (top-left) and on β-ta (top-right); ω-2θ maps around the Cu(111) peak for a copper film grown on α-ta (bottom-left) and on β-ta (bottom-right). Figure 5. Evolution with time of the Cu(111) peak intensity (top) and resistivity (bottom) in a 70nm Cu films. The selection of the barrier stack might in general influence the structural properties of the ECD copper overlayer, which in turn are related to its conductivity and electromigration properties. An example of how the structural properties of the barrier layer affect the copper structure is shown in Figure 4, where the texture analysis of copper films with the same thickness, but grown on two Ta layers having different crystalline phase (α-ta and β-ta, respectively) are compared [2]. The copper layer grown on β-ta results more textured in the (111) direction than the film grown on TaNTa bilayer, where the Ta is in the α-phase, as revealed by the higher intensity of the (311) peak obtained at low chi. Although a more detailed analysis of the orientation distribution function should be performed, our texture analysis (optimized for Cu), reveals that the most intense peak of β-ta obtained at low chi is the one related to the (002) orientation, whereas for TaNTa, the (110) and the (211) peak of α-ta (appearing at very different angles) are obtained simultaneously, confirming the more random polycrystalline nature of TaNTa as compared to Ta grown directly on the dielectric layer. The different grade of the structural order of copper, depending on the phase of Ta, is confirmed by the ω-2θ maps around the Cu(111) peak, reported in Figure 4. Although the experimental conditions were slightly different, and considering that the thickness of the two copper layers is the same, these maps confirm that Cu is more (111) oriented when it is deposited on β-ta. The Cu(111) orientation is generally desired, as it is usually associated to a higher electrical conductance. Using XRD and sheet resistance measurements, we confirmed this relationship and we also showed 72

74 that the structural properties of the copper stack are not stabilized after the film is deposited. This phenomenon, usually referred to as aging or selfannealing, was studied by combining XRD analysis and sheet resistance measurements as a function of time [5]. The relevant results for a 70 nm ECD copper layer deposited on a 25 nm PVD TaNTa barrier stack are reported in Figure 5. An increase with time of the Cu(111) component is clearly detected, together with the simultaneous decrease of film resistivity. The copper film stabilizes after 200 hours of storage in ambient air. [1] R. Rosenberg, D. C. Edelstein, C.-K. Hu, and K. P. Rodbell, Annu. Rev. Mater. Sci (2000). [2] G. Tallarida, C. Wiemer, L. Aina, S. Alberici, E. Ravizza, A. Giussani, G. Pavia, E. Varesi, MDM-D [3] G. Tallarida, C. Wiemer, S. G. Alberici, G. Brunoldi, S. Guerrieri, MDM-D [4] G. Tallarida, C. Wiemer, S. Guerrieri, S. Grasso, E. Ravizza, S. Spadoni, MDM-D [5] G. Brunoldi, S. Guerrieri, S. G. Alberici, E. Ravizza, G. Tallarida, C. Wiemer, T. Marangon, Microelectronic Engineering 82 (2005)

75 2. Emerging Materials for nanoscale CMOS devices 2.1 Rare earth oxides on Si for logic and memory applications: binary and ternary compounds 2.2 Band alignment at the La 2 Hf 2 O 7 /(001)Si interface 2.3 Dielectrics for channel materials alternative to Si: towards strained-si, Ge and GaAs devices 2.4 GeO 2 films grown on Ge subtrates by Atomic Layer Deposition and Molecular Beam Epitaxy 2.5 HfO 2 as gate dielectrics for Ge-based devices 2.6 Epitaxial HfO 2 on high-mobility semiconductors: theory and experiment 2.7 Epitaxial Gd 2 films on Ge The continuous scaling down of planar CMOS devices, in a medium-long term scenario (from the 45 nm node beyond the 22 nm node), requires the introduction and integration of new materials and architectures. The 2005 International Technology Roadmap for Semiconductors (ITRS- clearly highlights the challenges and potential solutions to materials limited device scaling. During the next ten years, front-end processes for MOSFET gate stack, as well as for DRAM and flash-memory storage devices, will require the introduction of a variety of high-κ materials as gate or inter-poly dielectrics, a precise control of oxide/semiconductor interfaces, highly-engineered metal films, and the incorporation of mobility-enhanced channels. Beyond the 22 nm node, it is still controversial whether planar CMOS will be a practical option. At that point, new device structures, such as FinFET or double-gate MOSFETs, as well as nanotubes- or nanowiresbased devices might be introduced in order to meet performance requirements. However, these devices have their own limitations in terms of mass production. In this chapter, we describe the efforts of MDM in addressing some of these issues. Long-term solutions for the scaling of gate and inter-poly dielectrics require the identification of materials with a high dielectric constant (>20) with outstanding electrical characteristics in terms of interface-state densities, reliability and charge retention. Therefore, beyond the 45 nm technology node, further downscaling might require the substitution of Al 2 and Hf-based dielectrics (chapter 1) with group III, rare earth oxides and ternary oxides. MDM has been pursuing an intense activity in the growth and characterization of binary and ternary rare earth compounds by ALD, as presented in the paragraph 2.1. In order to guarantee the requested gate leakage specification and the reliability requirements, the high-κ dielectrics must have a band gap of at least 4 5 ev and hole/electron barrier height >1 ev. These issues are addressed more into details in paragraph 2.2 for the LHO/Si system both from an experimental and theoretical point of view. 74

76 A very difficult challenge in device scaling is channel mobility enhancement, obtained as a first choice using strained-si channel and after substituting the Si channel with new materials. In order to introduce into planar CMOS devices alternative channel materials, the challenge is to find a proper gate stack providing also a good passivation of the oxide/semiconductor interface, low EOT and reasonable band alignment. An overview of the MDM activity for the development and characterization of dielectrics for strained- Si, Ge and GaAs channel is presented in paragraph 2.3, where the deposition, structural/chemical characterization as well as electrical properties and band alignment are discussed. Among the semiconducting materials, Ge offers a great potential for the integration into Si CMOS technology since it has bulk hole mobility four time higher than Si. Nevertheless, the unstable Ge oxide with the consequent easy formation of Ge suboxides (GeO x ) and high-density of interface states has been a serious point stopping the Ge technology. The growth of stoichiometric and well controlled GeO 2 films on Ge substrate, as well as Ge oxidation issues, are addressed in paragraph 2.4. Recently, the research on Ge-MOSFET has focused worldwide in the development of a reliable gate oxide, such as germanium oxynitride, Al 2, ZrO 2, HfO 2. MDM is also very active in this field investigating several gate dielectrics grown by ALD and MBE, such as rare earth oxides (paragraph 2.3), HfO 2 and Gd 2. HfO 2 is definitely the most studied and promising dielectric for Ge-based devices. Paragraph 2.5 presents a complete characterization of HfO 2 films grown by ALD and MBE on Ge, with focus on their structural and electrical properties, band alignment and Ge diffusion into the film. Finally, for a future aggressive scaling of EOT, epitaxial oxides which can avoid the formation of an interfacial layer with low-κ value, can be very promising. The growth and characterization of epitaxial HfO 2 and Gd 2 is discussed, respectively, in paragraph 2.6 and 2.7. STMicroelectronics (high-κ materials for non-volatile memory devices) as well as in the framework of several national and international research projects. The activity on rare earth oxides and rare earth based compounds is partially supported by the EU FP6 project REALISE ( ), the PAIS- INFM REOHK, and the joint Project between Italy and Russia funded by the Italian Ministry for Foreign Affairs ( ). Epitaxial and amorphous oxides (Hfbased and rare earth) grown by MBE are investigated in the EU FP5 STREP project INVEST ( ). The activity related to gate oxides for alternative channel materials, such as Ge and GaAs, are partially supported by the EU FP6 IST project ET4US ( ), the MAE Italia-Poland joint research project (2005) and the MAE Italia-Austria joint research project ( ). The activities presented in this chapter are carried out in close collaboration with 75

77 2.1 Rare earth oxides on Si for logic and memory applications: binary and ternary compounds G. Scarel 1, C. Wiemer 1, S. Spiga 1, E. Bonera 1, G. Seguini 1, G. Tallarida 1, X. Li 1, M. Fanciulli 1, Y. Lebedinskii 2, A. Zenkevich 2, I.L. Fedushkin 3, V. Fiorentini 4, F. Boscherini 5, S.D. Elliott 6, S. Schamm 7, G. Pavia CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy Moscow Engineering Physics Institute, Moscow, Russian Federation G. A. Razuvaev Institute of Organometallic Chemistry, Russian Academy of Sciences, Nizhny Novgorod, Russia INFM-SLACS, Sardinian Laboratory for Computational materials Science, and Department of Physics, University of Cagliari, Monserrato, Italy Department of Physics, University of Bologna, Italy Tyndall National Institute, Cork, Ireland CEMES-CNRS, Toulouse, France STMicroelectronics, Agrate Brianza, Italy Rare earth (RE) oxide films at MDM were deposited using atomic layer deposition [1,2,3]. Cyclopentadienyltype of precursors were shown to be very efficient and promising, i.e. the correct stoichiometry was easily achieved in the thin films [1,2,3]. For Lu, a new biscyclopentadienyl-type of precursor was syntesized [4]: [(η 5 -C 5 H 4 SiMe 3 ) 2 LuC1] 2. RE-based compounds were chosen for investigation at MDM mainly because the properties of the RE elements generate compounds with various characteristics. Among them the most interesting ones can be chosen, according to the needs of advanced microelectronics. The RE elements are the 15 elements of the Periodic Table (La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb, and Lu) with atomic numbers from 57 through 71. Among them, Pm is radioactive and does not occur naturally, but might be prepared synthetically. In the outer electronic configuration of the RE element row, the 6 s 2 shell is always occupied, the 5 d 1 configuration appears in La, Ce, Gd and Lu, and finally the 4 f shell is progressively filled as the atomic number increases. The degree of filling of the 4 f shell is therefore the distinctive characteristic of the RE elements. In particular, the half-filled (Gd, with 7 electrons in the 4 f shell) and the totally filled (Lu, with 14 electrons in the 4 f shell) configurations seem particularly stable. In the solid state, all 15 RE elements have the oxidation state +3, however some are stable also in the oxidation state +4 (Ce, Pr, and Tb), and other in the oxidation state +2 (Sm, Eu, Tm, and Yb). The classification of di-, tri-, and tetravalent RE elements and the RE-O bond lengths are summarized in Ref. [5]. It is noteworthy that the +4 oxidation state appears in elements that follow one with a stable configuration, whereas the +2 oxidation state appears in elements that precede one with a stable configuration. The phase of RE oxides depends on the RE atomic number, and so does the value of the dielectric constant (κ) of the corresponding oxides. Large κ values can be achieved in oxides crystallized in the hexagonal P6 3 /mmc phase [3,6]. This phase is however difficult to stabilize for oxides of RE elements other than La: generally the cubic bixbyite Ia3 phase is detected in RE oxides [1]. The oxides crystallized in the latter form offer a low κ ( 11-12) value mainly because of the relatively high energy of the transverse optical phonon with minimum frequency (Figure 1) [1,6,7]. Experiments carried out in order to manipulate the interfacial layer (IL) between the RE oxide and Si(100) indicated that there are processing methods to eliminate, or, at least, to significantly decrease in thickness, the IL [8]. These procedures promote dielectric stacks in real metaloxide-silicon (MOS) capacitors with higher κ than φ Δω φ φ Δω φ Δω Figure 1. (a)-(d) Transmission infrared spectra of cubic bixbyite Lu 2 on Si (solid lines). (a)(c)(d) Experimental spectra for various incidence angles φ. (b) Normal transmission infrared spectrum from density functional theory. All the experiments were in the far-infrared configuration. The resolution Δω is shown on the left of each spectrum. The dotted lines are the spectral simulations. 76

78 Lu). Indeed, these elements in films more than 10 nm thick are found to promote oxides with low sensitivity to moisture, and slight tendency to incorporate Si from Si(100) upon annealing [2,11,12]. Figure 2. Fundamental energy gap of the RE sesquioxides. Optical gaps are shown as circles connected by a dashed line, while the band gap values deduced from high temperature conductivity experiments are shown in diamonds. The origin of the small values of the conductivity gaps is not known. those with IL really existing between the RE oxide and Si(100). The absence of IL might lower the capacitance equivalent oxide thickness (CET) of the dielectric layer [7]. Because of the electronic configuration issues related to RE elements described above, there are periodic trends in the RE oxides. The latter concern e.g. the electronic properties of RE oxides. In particular, the band gap is largest for the oxides of the RE elements with empty (La), half-filled (Gd), and totally filled (Lu) 4 f shell (Figure 2) [9]. We argue that also conduction band offset (CBO) on Si(100) undergoes a periodic trend on Si(100) with maxima found for RE oxides with the highest band gap [5,9,10]. Other properties of RE elements change monotonously, such as those depending on the ionic radii of the RE elements. The RE atomic radii indeed decrease as the atomic number increases this is the so-called lanthanide contraction - from nm in La to nm in Lu [5]. The lanthanide contraction is due to the increase in localization of the 4 f shell as the atomic number increases. This phenomenon has dramatic consequences on the moisture sensitivity of RE oxides, and on the ability of Si from Si(100) substrates to diffuse into RE oxide layers upon annealing. E.g. La 2 films are extremely hygroscopic [3], while Yb 2 and Lu 2 ones are not [1,2]. Similarly, Si diffuses in a significant amount from Si(100) to La 2 films upon annealing [3], while Yb 2 and Lu 2 layers are less sensitive to this problem [1,2]. An exception to the general rule just described occurs for very thin films: the latter appear mostly in the hydroxide form and are Si-rich, even if the RE element has a small ionic radius (i.e. Yb and Thin Yb 2 and Lu 2 films on both p- and n-si(100) with well-shaped capacitance-voltage characteristics without significant dispersion of the accumulation capacitance were prepared [7]. Nevertheless, negative shift of the flat band voltage (indicating high density of fixed positive charges), and hysteresis loops were also detected [7]. In addition, the measured leakage currents for films with CET < 3 nm are not significantly lower than those measured in SiO 2 stacks of similar CET. Leakage currents are enhanced in crystalline films, as grown or after annealing. Therefore, in order to decrease these currents, ternary oxides are often preferred over binary ones because they are amorphous up to high temperatures ( 1000 C), and also because they promise κ values in some cases even higher than those of binary oxides. RE aluminates and silicates are among the ternary compounds that could be of interest for applications in advanced microelectronics. Indeed, La aluminate films deposited using Lacyclopentadienyl precursors [13] exhibit promising electrical characteristics, with κ = 20±3 (Figure 3) [14]. On the other hand, Lu silicates obtained from Si-containing Lu precursors ([(Me 3 Si) 2 N] 3 Lu [15], or [(η 5 -C 5 H 4 SiMe 3 ) 2 LuC1] 2 [16]), exhibit also interesting electrical properties [15], but also a low stack κ value ( 5 7). Figure 3. C-V curves for a 22 nm thick La aluminate film acquired at various frequencies. Inset: Capacitance equivalent oxide thickness (CET) values as a function of the La aluminate thickness. Gate area: (7.7±0.3) x 10-4 cm 2. Summarizing, RE oxides and RE-based compounds are potentially interesting for applications in Si(100)- based microelectronics. The κ values for binary oxide films crystallized in the cubic bixbyite structure, which 77

79 range between 11 and about 15, are moderately interesting for ultra-advanced applications on Si(100). Therefore, significant efforts are devoted at MDM to explore two alternative solutions to the problem of these low κ values. One solution consists in the stabilization of specific crystallographic phases, such us the hexagonal phase of La 2, which are capable to increase significantly the oxide κ value. Applications based on hexagonal La 2 films might be however limited due to the high oxide physical thickness needed to establish the interesting crystalline phase. Suitable choice of RE precursors (in particular modified La cyclopentadienyl-based precursors) might overcome this problem. Intensive research is going on in this direction. Another solution consists in the development of RE-based ternary compounds with κ values 20, and with thermodynamical stability on Si(100) for applications devoted to post-45 nm node devices. Promising results were already obtained for LaAl films. Finally, RE-based compound can be promising also for ultra-scaled devices based on channels alternative to Si, such us Ge and GaAs, as outlined in paragraphs 2.3 and 2.7. This work was supported by various Projects: Industrial MDM- STMicroelectronics project, European Project FP6 REALISE, Joint Project between Russia and Italy funded by the Russian Government (2005), PAIS-INFM REOHK, Joint Project between Italy and Russia funded by the Italian Ministry for Foreign Affairs ( ). Mr. M. Alia (CNR-INFM MDM National Laboratory) is acknowledged for sample preparation and processing. [1] G. Scarel, E. Bonera, C. Wiemer, G. Tallarida, S. Spiga, M. Fanciulli, I.L. Fedushkin, H. Schumann, Y. Lebedinskii, and A. Zenkevich, Appl. Phys. Lett. 85, 630 (2004). [2] M. Malvestuto, G. Scarel, C. Wiemer, M. Fanciulli, F. D Acapito, and F. Boscherini, Nucl. Instrum. Meth. in Phys. Res. B 246, 90 (2006). [3] X.L. Li et al., unpublished. [4] H. Schumann, I.L. Fedushkin, M. Hummert, G. Scarel, E. Bonera, and M. Fanciulli, Z. Naturforsch. 59b, 1035 (2004). [5] G. Scarel, A. Svane, and M. Fanciulli, in Rare earth oxide thin films: growth, characterization, and applications, Eds. M. Fanciulli and G. Scarel, Topics in Applied Physics 106, p. 1, Springer-Verlag (2007). [6] E. Bonera, G. Scarel, M. Fanciulli, P. Delugas, and V. Fiorentini, Phys. Rev. Lett. 94, (2005). [7] S. Spiga, C. Wiemer, G. Scarel, O. Costa, and M. Fanciulli, in Rare earth oxide thin films: growth, characterization, and applications, Eds. M. Fanciulli and G. Scarel, Topics in Applied Physics 106, p. 203, Springer-Verlag (2007). [8] S. Schamm, G. Scarel, and M. Fanciulli, in Rare earth oxide thin films: growth, characterization, and applications, Eds. M. Fanciulli and G. Scarel, Topics in Applied Physics 106, p. 153, Springer-Verlag (2007). 2.2 Band alignment at the La 2 Hf 2 O 7 /(001) Si interface G. Seguini 1, S. Spiga 1, E. Bonera 1, M. Fanciulli 1, A. Reyes Huamantinco 2, C.J. Först 2, C.R. Ashman 2, P.E. Blöchl 2, A. Dimoulas 3, G. Mavrou [9] G. Seguini, E. Bonera, S. Spiga, G. Scarel, and M. Fanciulli, Appl. Phys. Lett. 85, 5316 (2004). [10] M. Malvestuto, M. Pedio, S. Nannarone, G. Pavia, G. Scarel, M. Fanciulli, and F. Boscherini, J. Appl. Phys. 101, (2007). [11] Yu. Lebedinskii, A. Zenkevich, G. Scarel, and M. Fanciulli, in Rare earth oxide thin films: growth, characterization, and applications, Eds. M. Fanciulli and G. Scarel, Topics in Applied Physics 106, p. 127, Springer-Verlag (2007). [12] A. Zenkevich, Yu. Lebedinskii, G. Scarel, and M. Fanciulli, in Defects in High-κ Gate Dielectric Stacks. NATO Science Series. II. Mathematics, Physics and Chemistry 220. Ed. E. Gusev, p. 147, Springer-Verlag (2006). [13] G. Scarel, C. Wiemer, M. Fanciulli, A. Zenkevich, and Y. Lebedinskii, Internal report MDM-STMicroelectronics, D2, October [14] S. Spiga, G. Scarel, C. Wiemer, and M. Fanciulli, Electrical properties of La aluminate films grown by ALD, Internal report MDM-D [15] G. Scarel, C. Wiemer, G. Tallarida, S. Spiga, G. Seguini, E. Bonera, M. Fanciulli, Y. Lebedinskii, A. Zenkevich, G. Pavia, I.L. Fedushkin, G.K. Fukin, and G.A. Domrachev, J. Electrochem. Soc. 153, F271 (2006). [16] S.D. Elliott, G. Scarel, C. Wiemer, M. Fanciulli, Y. Lebedinskii, A. Zenkevich, and I.L. Fedushkin, in 15 th European Conference on Chemical Vapor Deposition, EUROCVD15. Eds. A. Devi, R. Fischer, H. Parala, M.D. Allendorf, and M. Hitchman. PV , 605, The Electrochemical Society, Pennington NJ (2005). CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy Institute for Theoretical Physics, Clausthal University of Technology, Clausthal-Zellerfeld, Germany Institute of Materials Science, NCSR Demokritos, Athens, Greece In order to integrate high-dielectric constant (highκ) gate oxides in future ultra-scaled devices, a large bandgap and as well as large (> 1 ev) conduction (CBO) and valence (VBO) band offsets, are key requirements. In this contribution we determine experimentally the 78

80 band alignment and the band gap of La 2 Hf 2 O 7 (LHO)- based metal-insulator-semiconductor (MIS) structures by internal photoemission (IPE) spectroscopy, photoconductivity (PC) measurements, and optical absorption spectroscopy [1]. Moreover, density functional theory (DFT) calculations are performed to determine the VBO [1]. LHO amorphous films are grown by molecular beam deposition (MBD) on a (2x1) reconstructed (001) Si surfaces at 600 C by co-evaporating Hf and La metals in the presence of atomic oxygen beam [1]. A sample is also grown on a quartz substrate in similar conditions, in order to perform optical absorption measurements. measuring the transition of holes from the conduction band of Si to the valence band of LHO oxide. The extrapolated value of the barrier height φ h at zero field is 3.5 ev. Considering the Si gap energy, the extracted value of the VBO is 2.4 ev. The photoconductivity response allows the evaluation of the oxide transport band gap (E G ). We found E G =5.6±0.1 ev (Figure 2(a)). The optical absorption coefficient α plotted in Figure 2(b) shows a square dependence on the photon energy which evidences the direct nature of the process. The 5.7±0.1 ev falls in the region of the onset of change in photoconductivity evidenced in Figure 2(a). The photocurrent was measured in the Al/LHO/n-Si MIS capacitors biased at different voltages and under monochromatic light irradiation (photon spectral range hν= ev). The IPE quantum yield Y is extracted by normalizing the photocurrent with respect to the incident photon flux. Figure 2. (a) Photoconductivity spectrum for the Al/LHO/n-(001)Si stack for a 2.5 V bias. (b) Optical absorption coefficient for LHO grown on quartz. Figure 1. (a) Internal photoemission spectrum of Al/LHO/n-(001)Si for a +1.8 V bias. (b) Internal photoemission spectrum for a -1.0 V bias. Figure 1(a) shows the behavior of Y 1/3 as a function of photon energy for gate bias of +1.8 V. This behavior is typical of a semiconducting emitter, Si in this case. A clear spectral threshold is identified at 2.9 ev. The extrapolation of the spectral thresholds, as a function of the applied fields, positions the zero-field threshold at 3.2 ev. This corresponds to the barrier height φ e between the valence band of Si and the conduction band of the oxide. The CBO is therefore equal to 2.1 ev, obtained by subtracting the Si gap (1.1 ev). By performing IPE on negatively biased MIS it is possible to probe the energy barrier between metal and LHO (Figure 1(b)). The Al/LHO barrier height is determined to be 2.8 ev. The same spectrum allows While the experimental results are obtained from an amorphous sample, the theoretical DFT calculations presented in this work are carried out on a crystalline system with cube-on-cube epitaxy of the metal oxide on Si(001). DFT calculations [1] have been performed for the crystalline LHO/Si(001) interface using the Projector Augmented Wave (PAW) method [1]. The VBOs are obtained from multilayer (sandwich) calculations in a way analogous to the experimental determination from x-ray photoelectron spectroscopy. The crystalline interfaces have been constructed by matching tetragonal LHO to Si(001), with an alignment of (001) LHO //(001) Si and [110] LHO //[110] Si orientations. We consider only interfaces that are electronically saturated and therefore do not pin the Fermi level. Deviations from the La/Hf bulk stoichiometry are necessary for electronically inactive interfaces. By changing the stoichiometry of the interfacial cation and anion layers from La 2 Hf 2 O 7 to La 3 HfO 5 or to LaHf 3 O 6 we can avoid the net charge at the interface and Fermi-level pinning. In the following, we investigate one class of interfaces, namely the La 3 HfO 5. Two of the three resulting interfaces having low energy are shown 79

81 Figure 3. Atomic structures of the simulated interfaces between LHO and Si(001). Top: Interface A. Bottom: Interface B. Interface A differs from the interface B by the orientation of the Si crystal relative to that of the LHO. This work was supported by the European Project INVEST. [1] G. Seguini, S. Spiga, E. Bonera, M. Fanciulli, A. R. Huamantinco, C. J. Först, C. R. Ashman, P. E. Blöchl, A. Dimoulas, and G. Mavrou, Appl. Phys. Lett. 88, (2006). 2.3 Dielectrics for channel materials alternative to Si: towards strained-si, Ge and GaAs devices S. Spiga 1, C. Wiemer 1, G. Scarel 1, G. Seguini 1, M. Perego 1, S. Ferrari 1, S. Baldovino 1, M. Fanciulli 1, A. Zenkevich 2, Y. Lebedinskii 2, Y. Panayiotatos 3, A. Dimoulas 3 in Figure 3. Within the accuracy of our calculations, all three interfaces can be considered degenerate. They differ in the orientation of the Si substrate relative to the oxide and/or the final position of the O atom that does not bind to Si. In interface A (Figure 3 top), the most stable interface, four O atoms remain at the interface saturating the Si dangling bonds, while the fifth O atom undergoes a large relaxation of 1.88 Å from the interface into the oxide, creating an electric dipole perpendicular to the interface. In interface B (Figure 3, bottom) all five O atoms remain at the interface in bulk-like positions. For interface B, the orientation of the Si substrate is perpendicular to that in interface A. From the relative positions of the valence band maxima of LHO and Si, we obtain VBOs of 1.81±0.27 and 2.37±0.02 ev for interfaces A and B, respectively. The large difference is due to the different position of the fifth O atom. Its relaxation induces an electric dipole responsible for the change in the band offsets. In conclusion, we have shown that, from the point of view of the band structure, a MIS structure based on MBE-deposited LHO is compatible with the future requirements of microelectronics. Theoretical arguments indicate that the local stoichiometry at the interface has to differ from the bulk to avoid Fermi level pinning in the conduction band. The calculations indicate that a variety of different structures coexist in a similar energy window. The calculated VBO values vary within a range of 0.5 ev and overlaps with the experimental data for the amorphous system CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy Moscow Engineering Physics Institute, Moscow, Russian Federation Institute of Materials Science, NCSR Demokritos, Athens, Greece Future high-performance devices will require the substitution of the silicon channel with high-mobility semiconductors, such as strained-si, Si-Ge, Ge and GaAs. Even if the optimization of the gate oxide for these substrates is still in the early stages, Ge- and GaAs-based metal oxide semiconductor (MOS) transistors incorporating high-κ dielectrics have been already successfully fabricated [1,2]. Al 2 and HfO 2 dielectrics have been largely studied for Si-based devices and considered as a first choice also for high-mobility channels. For instance, Al 2 deposited by ALD provides a good passivation for GaAs and InGaAs semiconductors [2]. However, the Al 2 dielectric constant is not high enough for ultra-scaled devices. HfO 2 (paragraphs 2.5 and 2.6) is therefore investigated as gate oxide, and recently, rare earth based dielectrics are gaining increasing interest. Besides the Gd 2 and (Ga 2 ) 1- (Gd O ) systems, which were deeply studied x 2 3 x for GaAs [3,4], other binary/ternary rare earth compounds, such as La 2, LaAl, Lu 2, CeO 2 are currently under evaluation for Ge and GaAsbased devices [5-8]. 80

82 In this contribution we report on the development of advanced gate oxides for high-mobility semiconductors, and on their structural, chemical and electrical properties, including the oxide/ semiconductor band alignment. Various high-κ oxides are deposited on strained Si, Ge and GaAs using atomic layer deposition (ALD) and molecular beam deposition (MBD). ALD technique provides uniform and conformal films, moreover the choice of precursor combination and growth parameters can be used to tailor the properties of the high-κ oxide stacks. On the other hand, MBD technique provides an additional capability for the in-situ cleaning/preparation of the semiconductor surfaces and the in-situ deposition of proper passivation layer, such us GeON on Ge [9] and Si on GaAs [10]. In the following the results obtained for Al 2 deposited on strained Si and Ge and for a selection of rare earth oxides grown by ALD and MBD on Ge, and GaAs, are presented. Figure 1. Interface traps density measured using the Hill-Coleman method for Al/10 nm Al 2 /Si-strained devices as a function of increasing strain in the channel (which is related to the Ge content in the SiGe buffer layer under the Si-strained channel). Al 2 films were deposited at 300 C on strained- Si/SiGe/Si stacks (with variable Ge content in the SiGe layers) and on bulk Ge [6] using the thrimethylaluminum Al(CH 3 ) 3 or TMA and H 2 O precursors. MOS capacitors were fabricated by Al metal evaporation through a shadow mask to test electrical properties. Al 2 films are amorphous and with a sharp interface either on Ge and on strained- Si channel. Al 2 films exhibit good electrical properties on both Ge and strained-si substrates, with low leakage currents, even if for Ge capacitors the interface trap density (D it ) is in the ev -1 cm -2 range [6]. Moreover for the strained-si channel, an interesting result is that the increase of strain is effective in reducing interface trap density as evidenced by conductance measurements (Figure 1) and electrically detected magnetic resonance results (not shown) [11]. Our results are in agreement with those reported in the literature [12]. Lu 2 films were deposited on Ge and GaAs at 360 C using the [(η 5 -C 5 H 4 SiMe 3 ) 2 LuCl] 2 as Lu source and H 2 O as oxygen source. Yb 2 films were grown on Ge at 360 C using two precursor combinations: Yb(C 5 H 5 ) 3 + H 2 O or Yb(C 11 H 19 O 2 ) 3 +. Lu 2 films deposited on Ge and GaAs are nanocrystalline, with grain size always lower than film thickness. The crystallites exhibit the cubic bixbyite structure in the Ia-3 space group (Figure 2, top). The electronic density measured by x-ray reflectivity (XRR) is 2.03±0.05 e - /Å 3 (the nominal electronic density of Lu 2 is 2.18 e - /Å 3 ). The Lu 2 /Ge and Lu 2 /GaAs interfaces are sharp, and both XPS and XRR (Figure 2, bottom) techniques do Figure 2. XRD data (top) and XRR spectra and corresponding fittings (lines), (bottom) of Lu 2 films: 13.6 nm thick film on Ge(squares) and 10.7 nm thick film on GaAs (circles). not evidence the formation of any interfacial layer (IL) [8]. Yb 2 films in the 5-30 nm range deposited on Ge using both precursor combinations are crystallized as grown in the cubic bixbyite structure. Figure 3 shows 81

83 of CeO 2 (Figure 4). Moreover, a thick Ge-O-Ce interfacial layer (IL) is spontaneously formed during the growth, as revealed by XPS and cross-sectional TEM. The CeO 2 films exhibit good MOS behaviour (Figure 5) demonstrating that an oxidized Ge IL could be beneficial for the electrical properties in terms of leakage current and D it [5]. On the other hand, this IL increases the equivalent oxide thickness and its thickness must be reduced for device scaling. Figure 4. Grazing incidence XRD spectra of 15 nm thick CeO 2 films on Ge grown by MBD at 225 C. In the inset the cros-sectional TEM image (courtesy of M. Seo, EPFL, Switzerland). Figure 3. Top. CV and JV (inset) characteristics of Al/Lu 2 /n-ge capacitor. The Lu 2 film is 14 nm thick. Bottom: CV characteristics of Al/10 nm Yb 2 /n-ge capacitor. The Yb 2 film is deposited using the Yb(C 11 H 19 O 2 ) 3 and precursors. Gate area: 7.8 x10-4 cm 2. two representative CV characteristics acquired for Al/ Lu 2 /Ge and Al/Yb 2 /Ge capacitors. Film thickness is in the nm range. Despite the relatively low dielectric constant values measured for Lu 2 and Yb 2 (in the range), these rare earth dielectrics exhibit reasonably good CV characteristics on Ge without the need of any intentional interfacial layer, such as GeON which is necessary for HfO 2 and LaAl oxides on Ge [5,9]. Moreover, the use of ozone as oxygen precursor further improves the Yb 2 /Ge interface quality reducing the frequency dispersion of the CV curves (Figure 3, bottom) and the lowering D it down to the ev -1 cm -2 range. The CV curves of Yb 2 films grown using H 2 O as oxygen source (not shown) are similar to those of the Lu 2 /Ge stack, and the measured D it is in the ev -1 cm -2 range. CeO 2 films were deposited by MBD on clean Ge surfaces at various temperatures (from room temperature up to 360 C) [5]. Grazing incidence x- ray diffraction analyses revealed that the films (10-20 nm thick) are polycrystalline in the cubic phase Figure 5. CV characteristics acquired at various frequencies for Al/CeO 2 / Ge capacitors. Another among the most important requirements imposed to gate oxides in ultra-scaled devices is that they must have conduction band offset (CBO) and valence band offset (VBO) larger than at least 1.0 ev, in order to avoid high leakage current. X- ray photoelectron spectroscopy (XPS) and internal photoemission spectroscopy (IPE) are powerful techniques to characterize the band alignment of high-κ dielectrics on semiconductors [13-15]. The experimental CBO values reported for Al 2, ZrO 2 and 82

84 in developing new high-quality materials and the optimization of their interface with semiconductors. Rare earth binary oxides and their ternary compounds are promising candidates as gate dielectrics, providing that scalability and proper interface control are achieved. The activities presented in this paragraph are partially supported by the EU FP6 IST project ET4US and the MAE Italia-Poland joint research project. Figure 6. Top: IPE spectra of an Al/Al 2 /Ge stack for various positive applied voltages. Bottom: Energy band diagram for semiconductor-lutetia interface for different semiconductors (Si, Ge, GaAs). HfO 2 [16,17], Lu 2 [13,14] and ternary compounds (LaAl, LaSc, GdSc, DySc, La 2 Hf 2 O 7 ) [6,18] on Si are almost identical, ranging from 2.0 to 2.1 ev. Despite the increasing interest to integrate high-κ oxides in Ge- and GaAs-based devices, the band alignment has been measured only for a restricted number of oxides [6]. Figure 6, top, shows the cube root (Y 1/3 ) of IPE quantum yield versus photon energy measured for an Al/Al 2 /n-ge capacitor biased at positive voltages. The spectral thresholds, determined as a function of the applied voltage, are related to the electron transitions from the Ge valence band to the Al 2 conduction band. The extracted barrier energy at zero applied field (Φ e ) is equal to 2.8±0.1 ev. By subtracting the Ge band gap to Φ e, the CBO value is determined to be 2.1±0.1 ev. Similar measurements were performed for other oxides [6] grown on Si, Ge and GaAs [6,13,14,15]. Figure 6, bottom, shows, for instance, the results obtained for the band alignment of Lu 2 on Si, Ge and GaAs as measured using IPE. The results were confirmed by XPS. In summary, the search of suitable gate oxides for high-mobility channel devices requires a large effort [1] N. Wu, Q. Zhang, C. Zhu, D. S. H. Chan, A. Du, N. Balasubramanian, M. F. Li, A. Chin, J. K. O. Sin, and D.-L. Kwong, IEEE Electron Device Lett. 25, 631 (2004). [2] P. D. Ye, G. D. Wilk, B. Yang, J. Kwo, S. N. G. Chu, S. Nakahara, H.-J. L. Gossmann, J. P. Mannaerts, M. Hong, K. K. Ng, and J. Bude, Appl. Phys. Lett. 83, 180 (2003). [3] J. Kwo, D. W. Murphy, M. Hong, R. L. Opila, J. P. Mannaerts, A. M. Sergent, and R. L. Masaitis, Appl. Phys. Lett. 75, 1116 (1999). [4] M. Passlack, N. Medendorp, R. Gregory, and D. Braddock, Appl. Phys. Lett. 83, 5262 (2003). [5] Dimoulas, D. P. Brunco, S. Ferrari, J. W. Seo, Y. Panayiotatos, A. Sotiropoulos, T. Conard, M. Caymax, S. Spiga, M. Fanciulli, Ch. Dieker, E. K. Evangelou, S. Galata, M. Houssa, M. M. Heyns, Thin Solid Films (2007). [6] S. Spiga, C. Wiemer, G. Scarel, G. Seguini, M. Fanciulli, A. Zenkevich and Yu. Lebedinskii, in Advanced Gate Stacks on High Mobility Semiconductors, edited by A. Dimoulas, E. Gusev, P. McIntyre and M. Heyns, Springer-Verlag (2007). [7] A. Dimoulas, in: E. Gusev (Ed.), Defects in High-κ Gate Dielectrivc Stack s, NATO Science Series II, vol. 220, Springer, Netherlands, (2006), p.237. [8] Yu. Lebedinskii, A. Zenkevich, G. Scarel, and M. Fanciulli, in Rare earth oxide thin films: growth, characterization, and applications, Eds. M. Fanciulli and G. Scarel, Springer-Verlag (2006). [9] A. Dimoulas, G. Mavrou, G. Vellianitis, E. K. Evangelou, N. Boukos, M. Houssa, M. Caymax, Appl. Phys. Lett. 86, (2005). [10] S. J. Koester, E. W. Kiewra, Yanning Sun, D. A. Neumayer, J. A. Ott, M. Copel, D. K. Sadana, D. J. Webb, J. Fompeyrine, J.-P. Locquet, C. Marchiori, M. Sousa, and R. Germann, App. Phys. Lett. 89, (2006). [11] M. Fanciulli et al., unpublished (Oral presentation-mrs Fall Meeting, Boston (2006). [12] A. Stesmans, P. Somers, and V. V. Afanas ev, C. Claeys and E. Simoen, Appl. Phys. Lett. 89, (2006). [13] Seguini, M. Perego, M. Fanciulli, in Rare earth oxide thin film: growth, characterization, and applications, Eds. M. Fanciulli and G. Scarel, Springer-Verlag, (2007). [14] G. Seguini, E. Bonera, S. Spiga, G. Scarel, M. Fanciulli, Appl. Phys. Lett. 85, 5316 (2004). [15] M. Perego, G. Seguini, G. Scarel, M. Fanciulli, Surf. Interface Anal. 38, 494 (2006). [16] V. V. Afanas ev, M. Houssa, A. Stesmans, and M. M. Heyns. Appl. Phys. Lett. 78, 3073 (2001). [17] V. V. Afanas ev, A. Stesmans, F. Chen, X. Shi, and S. A. Campbell, Appl. Phys. Lett. 81, 1053 (2002). [18] V. V. Afanas ev, A. Stesmans, C. Zhao, M. Caymax, T. Heeg, J. Schubert, Y. Jia, D. G. Scholm, and G. Lucovsky, Appl. Phys. Lett. 85, 5917 (2004). 83

85 2.4 GeO 2 films grown on Ge substrates by Atomic Layer Deposition and Molecular Beam Epitaxy A. Molle 1, M. N. K. Bhuiyan 1, G. Tallarida 1, M. Perego 1, G. Scarel 1, M. Fanciulli 1, I. L. Fedushkin 2, A. A. Skatova 2 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy 12 G. A. Razuvaev Institute of Organometallic Chemistry, Russian Academy of Sciences, Nizhny Novgorod, Russia In Ge-based devices the formation of a passivation layer prior to (or during) the growth of the dielectric film is beneficial for improving the electrical performances of the Ge/dielectric stack [1]. Therefore, the study of Ge oxidation process is of fundamental importance to understand and improve the properties of the oxide/ semiconductor interface. Recently we demonstrated for the first time our capability to fabricate a pure and stoichiometric GeO 2 by atomic layer deposition (ALD) [2] and we also presented the use of atomic oxygen (AO) radicals as an alternative procedure to achieve an high efficiency on the oxidation process even at relatively low temperatures in a controlled ultra high vacuum (UHV) environment [3]. Here both growth methodologies were attempted and the resulting samples were investigated by means of X-ray photoelectron spectroscopy (XPS). GeO 2 layers have been grown via ALD by exploiting the solid compound 1,2-bis[(2,6- diisopropylphenyl)imino] acenaphthene germanium, [i.e. (dpp-bian)ge], as a valuable Ge source [4]. The divalent Ge compound (dpp-bian)ge, is a monomeric two-coordinate Ge derivative. Moreover, it is a solid precursor, user-friendly, easy to handle at room temperature, neither toxic nor flammable and volatile, as required for ALD. has been used both as oxygen source and as oxidation agent to transform the divalent Ge of the chemisorbed precursor into tetravalent Ge on the growing surface oxygen precursor. Before growth, the native oxide was removed by dipping n- type Ge(001) for 30 s in a dilute HF solution at room temperature (RT). The Ge substrate was then rinsed in de-ionized water for 30 s. The XPS scan of the Ge 3d region (not shown) for the Ge substrate after cleaning procedure indicates that the native oxide was completely removed leaving a clean Ge surface with very limited oxygen contaminations. According to the XPS analysis the deposited films are composed of perfectly stoichiometric GeO 2 with extremely low carbon contamination level and no suboxides at the GeO 2 /Ge interface. On the other hand, the formation of GeO 2 layers in ultra high vacuum (UHV) has been obtained by exposing n-type Ge(001) substrates to an AO beam at several substrate temperatures, ranging from RT to 400 C. The experiments were performed in a Molecular Beam Epitaxy (MBE) system (Omicron Nanotechnology GmbH) with a base pressure of mbar. The O 2 dissociation was performed by a radio-frequency atomic gas source. The Ge(001) surfaces was treated by several cycle of Ar + sputtering at RT and annealing at 500 C thus reaching an extremely high degree of structural order and chemical purity. The as-prepared Ge samples were promptly exposed to AO beam at a partial pressure of mbar for a fixed exposure time of 20 min, and varying the substrate temperature T ox. The Ge 3d peak can be deconvoluted into five components [5]. The related components in the XPS shape profile are the elemental peak (Ge B) at binding energy BE=29.6±0.2 ev and the oxidation state components, Ge 1+ (BE=30.3±0.2 ev), Ge 2+ (BE=31.4±0.1 ev), Ge 3+ (BE=32.4±0.1 ev) and Ge 4+ (BE=33.2±0.1 ev), that correspond respectively to the following oxide species: Ge 2 O, GeO, Ge 2 and GeO 2. Although the oxidation of the substrate takes place already at RT, the formation of a stoichiometric GeO 2 layer is remarkably enhanced when T ox =300 C, as deduced from the comparative diagram in Figure 1(a) in which the presence of a GeO 2 bonding configuration is identified by a chemical shift of 3.6 ev. At lower temperatures a significant concentration of sub-oxide species forms, whereas at T ox >300 C the surface chemical transformation of GeO 2 into sub-oxides sets in (see the XPS scan taken at T ox =400 C in Figure 1(a)). The different concentrations of the oxide species can be quantitatively determined from the multiple Voigt fitting to the XPS data. For clarity, while in the XPS Ge 3d scan taken for T ox =200 C (Figure 1(b)) various oxidation state components are indicative of a mixture of different oxide stoichiometries, the XPS scan taken at T ox =300 C presents a majority contribution due to the GeO 2 species as inferred from Figure 1(c). A GeO 2 /GeO x IL between Ge and high-κ oxides might modify the band structure, and, consequently, the electronic properties, of the entire stack. Even if the band alignment of a high-κ/ge heterojunction seems 84

86 (a) (b) Binding Energy (ev) (c) Binding Energy (ev) Figure 1. XPS Ge 3d spectra after 20 min long AO exposure at the substrate temperatures T ox = RT, 200 C, 300 C and 400 C (a); representative fit to the Ge 3d spectrum of the oxides formed at T ox =200 C (b) and at T ox =300 C (c). Ge GeO 2 CBM 0.6 ± 0.1 ev we carefully investigated the band structure of the GeO 2 /Ge heterojunction by XPS according to Kraut s method [8]. In Figure 2 we reconstructed the band structure of the GeO 2 /Ge heterojunction. A ~ 0.6 ev band bending was observed in the Ge substrate with GeO 2 on top. The band bending may originate from charged states at the GeO 2 /Ge interface and occurs in opposite directions for n-type and p-type Ge. Therefore different band offsets are expected to occur for GeO 2 /Ge heterojunction with n-type and p- type Ge substrates. In a very recent paper the VBO and CBO of GeO 2 on p-type Ge(001) were reported to be 4.0 ± 0.05 ev and 1.04 ± 0.05 ev respectively [9]. These values are consistent with the shift of the Ge 3d signal towards higher binding energy reported by Hovis et al. [10]. According to the data therein reported it is possible to calculate the VBO using Kraut s method [8]. The VBO value obtained in this way is perfectly in agreement with that measured by Otha et al. [9]. In conclusion we observed that the GeO 2 buffer layer introduces a high barrier for holes and a very low barrier for electrons. This result might have an incisive impact in the interpretation of the band alignment of a high-κ dielectric/geo 2 IL/Ge substrate configuration. This work has been partially supported by the European Project ET4US. CBM VBM 0.6 ± 0.1 ev VBM Figure 2. GeO 2 /Ge heterojunction band structure reconstruction according to XPS measurements. not to be modified by the formation of a GeO 2 -like IL [6], a GeO 2 -Ge interface was shown to introduce a barrier for holes in the band structure of the highκ/ge heterojunction [7]. Afanas ev et al. [7] showed that post deposition annealing treatment is effective in growing a GeO 2 -like buffer layer at the HfO 2 /Ge interface. According to the data therein reported, the presence of the IL does not affect electron current but introduces an additional sub-threshold for holes. In order to understand the effect of the GeO 2 buffer layer on the high-κ/semiconductor band structure, 4.5 ± 0.1 ev 5.81 ± 0.04 ev [1] S. Spiga, C. Wiemer, G. Tallarida, G. Scarel,S. Ferrari, G. Seguini, and M. Fanciulli, Appl. Phys. Lett. 87, (2005). [2] M. Perego, G. Scarel, M. Fanciulli, I. L. Fedushkin, and A. A. Skatova, Appl. Phys. Lett. 90, (2007). [3] A. Molle, M. N. K. Bhuiyan, G. Tallarida, and M. Fanciulli, Appl. Phys. Lett. 89, (2006); A. Molle, M. N. K. Bhuiyan, G. Tallarida, and M. Fanciulli, Mat. Sci. in Semicon. Proc. 9, (2006). [4] I. L. Fedushkin, A. A. Skatova, V. A. Chudakova, N. M. Khvoinova, a. Yu, Baurin, S. Dechert, M. Hummert, and H. Schumann, Organometallics 23, 3714 (2004). [5] D. Schmeisser, R. D. Schnell, A. Bogen, F. J. Himpsel, D. Rieger, G. Landgren, J. F. Morar, Surf. Sci. 172, 455 (1986). [6] M. Perego, G. Seguini, M. Fanciulli, J. Appl. Phys. 100, (2006). [7] V. V. Afanas ev, and A. Stesmans, Appl. Phys. Lett. 84, 2319 (2004). [8] E. A. Kraut, R. W. Grant, J. R. Waldrop and S. P. Kowalczyk, Phys. Rev. Lett. 44, 1620 (1980). [9] A. Ohta, H. Nakagawa, H. Murakami, S. Higashi, S. Miyazaki., J. Surf. Sci. Nanotech. 4, 174 (2006). [10] J. S. Hovis, R. J. Hamers, C. M. Greenlief, Surf. Sci. 440, L815 (1999). 85

87 2.5 HfO 2 as gate dielectrics for Ge-based devices S. Spiga 1, C. Wiemer 1, G. Scarel 1, G. Tallarida 1, G. Seguini 1, M. Perego 1, S. Ferrari 1, M. Fanciulli 1, G. Mavrou 2, A. Dimoulas 2, S. Kremmer 3, C. Teichert 3, G. Pavia CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy Institute of Materials Science, NCSR Demokritos, Athens, Greece Institute of Physics, University of Leoben, Austria STMicroelectronics, Agrate Brianza, Italy Figure 1. Grazing incidence XRD spectra of HfO 2 films deposited on Ge by ALD using HfCl 4 in combination with various precursors for the oxygen: H 2 O (black), Hf(O t Bu) 2 (mmp) 2 (red) and (green). Among the various high-κ materials, HfO 2 has been largely investigated for ultra-scaled Ge-based devices [1-15]. Despite several problems mainly related to the control of HfO 2 /Ge interface properties [1,9,14] and Ge diffusion into the dielectric [2,7,11,12 ], the fabrication of HfO 2 /Ge based transistors with good properties has been reported [5,10]. This contribution reports on the optimization of HfO 2 / Ge stacks grown using either atomic layer deposition (ALD) [1-4,8] and molecular beam deposition (MBD) [7,9,15]. Structural, chemical and electrical properties were characterized using a variety of techniques, with emphasis on the HfO 2 /Ge interface properties [1-4] and band alignment [1,3,8], Ge diffusion into the dielectric [7] and local electrical properties [15]. In the first part of the contribution we present the results obtained for ALD grown HfO 2 films, and in the second part those grown by MBD. HfO 2 films were grown by ALD at 375 C on (001) Ge substrates alternating pulses of HfCl 4 and, of HfCl 4 and H 2 O, or of HfCl 4 and Hf(O t Bu) 2 (mmp) 2 [1,3,4]. The latter compound was used as oxygen source alternative to H 2 O and. Ge wafers (Umicore) were cleaned in a diluted HF:H 2 O(1:50) solution and rinsed in deionized water. Grazing incidence X-ray diffraction (XRD) analyses (Figure 1) revealed that 10 nm thick HfO 2 films grown using H 2 O as oxygen source are crystallized in a mixture of monoclinic and orthorhombic phases. On the other hand films grown using exhibit only an early stage of crystallization, and those deposited using Hf(O t Bu) 2 (mmp) 2 are Figure 2. Cross-sectional TEM images of HfO 2 films grown by ALD - at 375 C using HfCl 4 and H 2 O. The distance among ten HfO 2 (111) planes is shown. mostly amorphous (only small and isolated grains are revealed by cross-sectional TEM). All the HfO 2 films exhibit an electronic density of 2.40±0.05 e - /Å 3 which is close to the value reported for bulk HfO 2 [16]. The various oxygen precursors strongly influence also the interface properties from the structural, chemical and electrical point of view. HfO 2 films deposited using H 2 O exhibit local epitaxial growth on Ge (Figure 2), as discussed more into details in paragraph 2.6. Films grown using Hf(O t Bu) 2 (mmp) 2 exhibit a sharp and smooth interface with Ge, while promotes the formation of a thick GeO x interlayer, as shown by cross-sectional TEM in Figure 3. Time of flight secondary ion mass spectrometry (ToF-SIMS) profiles show no significant Ge diffusion in the bulk of as grown HfO 2 films [1]. Carbon, hydrogen and chlorine contaminants are mainly accumulated close to the Ge interface. Their concentrations in films grown using are one (hydrogen and chlorine) or two (carbon) orders of magnitude lower than those measured in films grown using H 2 O and Hf(O t Bu) 2 (mmp) 2. 86

88 Figure 3. Cross-sectional TEM images of HfO 2 films grown by ALD at 375 C using HfCl 4 as metal source and Hf(O t Bu) 2 (mmp) 2 (a) or (b) as oxygen source. Metal oxide semiconductor (MOS) capacitors were fabricated by thermal evaporation of Al gates through a shadow mask; Al was deposited on the backside of the Ge wafers as ohmic contact. Figure 4(a) shows multi-frequency capacitance-voltage (CV) characteristics of a 9.6 nm thick HfO 2 film grown at 375 C on clean Ge surface. Significant frequency dispersion both in accumulation and in inversion is revealed. Sweeping the gate voltage from inversion to accumulation and back, a large hysteresis (> 400 mv) is detected. The same behavior was reported for ZrO 2 films grown on HF-last Ge at 300 C by ALD [17], and was attributed to a highly defective interface. The CV characteristics of HfO 2 films grown alternating pulses of Hf(O t Bu) 2 (mmp) 2 and HfCl 4 (not shown) exhibit exactly the same behaviour of HfO 2 films deposited from HfCl 4 and H 2 O. On the other hand, films deposited on Ge using and HfCl 4 exhibit a better MOS behaviour (Figure 4(b)). CV curves do not show any frequency dispersion in the inversion and accumulation regions in the 10 khz-300 khz frequency range. The D it estimated using the Hill- Colemann method is in the 0.5-1x10 12 cm -2 ev -1 range. These improved electrical properties using can be related to the GeO x interfacial layer which develops during the HfO 2 deposition. Capacitance equivalent oxide thickness (CET) values are extracted from the accumulation capacitance at 10 khz (without taking into account quantum mechanical corrections) for 3 to 10 nm thick films grown using. The HfO 2 dielectric constant is determined to be 17±1 from the slope of the linear fit of the CET values versus HfO 2 physical thickness (Figure 4(c)). The intercept with the y-axis gives a CET of 1.9 nm for the IL. Taking into account the IL physical thickness, its dielectric constant turns Figure 4. CV curves acquired at different frequencies for HfO 2 films grown on Ge by ALD using the HfCl 4 +H 2 O (a) or HfCl 4 + (b) precursors combinations. (c) CET as a function of HfO 2 physical thickness. out to be around 4.5, close to the one reported for GeO 2 (5-6) [18]. HfO 2 films were grown by MBD on differently prepared (001) Ge surfaces. The native oxide was desorbed in situ under ultra-high vacuum (UHV) conditions by heating the substrate at 360 C for 15 min until a (2x1) reconstruction appears in the reflection high-electron energy diffraction (RHEED) pattern, indicative of a clean (001)Ge surface. For selected samples, an ultrathin GeO x N y layer with variable N and O content was formed by exposing at 225 C the Ge surface to 87

89 atomic O and/or N beams generated by a rf plasma source with the simultaneous evaporation of Ge. Finally, 3-20 nm thick HfO 2 films were deposited at several temperatures between 60 and 360 C. The films structure (amorphous or polycrystalline) depends on the film thickness and on the deposition temperature [6,7,9]. Cross-sectional TEM and X-ray reflectivity data show a sharp HfO 2 /Ge interface for films deposited on a clean (2x1)-reconstructed Ge(001) surface. For the HfO 2 /GeO x N y /Ge stacks the detected interfacial GeO x N y layer is thinner than the nominal value (1 nm), due to the dissolution of the interfacial layer and the consequent Ge and/or GeO x diffusion into the HfO 2 films during the MBD growth process [7,12]. Figure 5 shows the ToF-SIMS depth profiles acquired for HfO 2 films deposited at 225 C on Ge (black curve) and GeO x N y /Ge (red curve). A significant amount of Ge is detected into the films grown on GeO x N y and GeO x (not shown), while for HfO 2 deposited directly on Ge there is no significant germanium diffusion. Similar ToF-SIMS profiles are acquired for all films deposited in the C temperature range. The phenomenon responsible for Ge incorporation into HfO 2 during the growth process is the reaction of the GeO 2 component present in the GeO x N y interlayer with the metallic Hf, leading to the formation/diffusion of the volatile GeO species. Despite the instability of the interfacial layers and Ge in-diffusion, the HfO 2 /GeO x N y capacitors exhibit good electrical properties with low leakage current and equivalent oxide thickness (EOT) less than 1 nm [6]. On the other hand, the HfO 2 /Ge stack exhibits poor electrical properties and high interface traps density. In order to understand better the role of the interfacial layer and its composition on the electrical properties, conducting atomic force microscopy Figure 6. Topography (left) and 2D current images (right) measured for 6.6 nm HfO 2 /GeO x /Ge and 7.8 nm HfO 2 /GeO x N y /Ge stacks. Figure 7. IPE spectra of the Al/HfO 2 /Ge samples grown using or H 2 O as oxygen precursors for various positive applied voltages. The lines represent the linear fittings. In the inset the photoconductivity part of the spectrum at a positive applied voltage of 1.1 V is depicted. : V= +1.0 V, V= +1.4 V, V= +1.8 V. H 2 O: V= +0.5 V, V= +0.7 V, V= +0.9 V. Figure 5.ToF-SIMS profiles for HfO 2 films grown at 225 C on Ge and GeON/Ge. The Hf and Ge signal are shown. Since the two films have different thickness, the ToF-SIMS profiles are aligned at the HfO 2 /Ge interface. Sputtering was performed using Cs + at 0.5 kev, while Ga + at 25 kev was used for the generation of the analyzed negative ions. (C-AFM) was used to investigate the electrical homogeneity of HfO 2 films. Measurements were performed in UHV using conductive diamond coated probes. Figure 6 shows the morphology (left) and the current distribution (right) images of HfO 2 /GeO x /n- Ge (top) and HfO 2 /GeO x N y /n-ge (bottom) stacks The voltage applied to the samples is 6 V. The distribution of leakage currents and surface morphology are collected simultaneously. The film grown on GeO x shows several leakage spots where the measured current is high and spread out on a large area. For the sample grown on GeO x N y, the leakage current events occur evenly all over the sample and are less severe than for HfO 2 /GeO x stacks, demonstrating that N is effective in improving the film electrical homogeneity [15]. 88

90 Finally, we measured the HfO 2 band alignment of films grown by ALD on Ge, and films grown by MBD on GeO x N y /Ge, using internal photoemission (IPE) and X-ray photoelectron spectroscopy (XPS) [1,8]. Figure 7 shows the cube root (Y 1/3 ) of IPE quantum yield versus photon energy measured for HfO 2 films grown by ALD using and H 2 O. The measured barrier energy is 2.7±0.1 ev in both cases. This barrier is related to the electron transition between the Ge valence band and the HfO 2 conduction band. By considering a Ge band gap of 0.7 ev the extracted conduction band offset (CBO) values are 2.0±0.1 ev in both cases. The photoconductivity spectrum (inset of Figure 7) gives a band gap value of 5.6±0.1 ev for the HfO 2 oxide grown either by H 2 O or by. The VBO values inferred by IPE data are 2.9±0.1 ev. VBO values of 3.0±0.1 ev and 3.1±0.1 ev were obtained by XPS for the samples grown using and H 2 O, respectively, in excellent agreement with those obtained by IPE and PC measurements [8]. The CBO measured for HfO 2 films grown by MBD on GeO x N y / Ge is 2.2±0.1eV, which is close to the value measured for HfO 2 films deposited by ALD. In summary, we presented the results obtained for HfO 2 films grown using ALD and MBD on Ge. We showed that the choice of precursor combination for the ALD growth allows to control structural and electrical properties and that HfO 2 /GeO x N y stacks grown by MBD exhibit sharp interface and good electrical properties with EOT < 1 nm. The HfO 2 band alignment was determined using IPE and XPS. The values of CBO and VBO are found to be independent, within experimental errors, from the oxide structure (amorphous, polycrystalline), the HfO 2 /Ge interface properties and deposition techniques. This work is partially supported by the European project ET4US (IST- FP6), and by the Ministry for Foreign Affairs in the framework of the scientific and technological collaboration between Italy and Austria (project 9/2004) and between Italy and Poland (2005). [1] S. Spiga, C. Wiemer, G. Tallarida, G. Scarel, S. Ferrari, G. Seguini, M. Fanciulli, Appl. Phys. Lett., 87, (2005). [2] M. Caymax, S. Van Elshotcht, M. Houssa, A. Delabie, T. Conard, M. Meuris, M.M. Heyns, A. Dimoulas, S.Spiga, M. Fanciulli, Mat. Sci. Eng. B, 135, 256 (2006). [3] S. Spiga, C. Wiemer, G. Scarel, G. Seguini, M. Fanciulli, A. Zenkevich and Yu. Lebedinskii, in Advanced Gate Stacks on High Mobility Semiconductors, edited by A. Dimoulas, E. Gusev, P. McIntyre and M. Heyns, Springer-Verlag (2007) in press (November 2007). [4] M. Fanciulli, S. Spiga, G. Scarel, C. Wiemer, G. Seguini, G. Tallarida, Mat. Res. Soc. Symp. Vol 786, E6.14 (2004). [5] N. Wu, Q. Zhang, C. Zhu, D.S.H. Chan, A. Du, N. Balasubamanian, M.F. Li, A. Chin, J. K. O. Sin, and D.-L. Kwong, IEEE Elctron Device Lett. 25, 631 (2004). [6] A. Dimoulas, G. Mavrou, G. Vellianitis, E. K. Evangelou, N. Boukos, M. Houssa, M. Caymax, Appl. Phys. Lett. 86, (2005). [7] S. Ferrari, S. Spiga, C. Wiemer, M. Fanciulli, A. Dimoulas, Appl. Phys. Lett. 89, (2006). [8] M. Perego, G. Seguini, M. Fanciulli, Energy band alignment of HfO 2 on Ge, J. Appl. Phys. 100, (2006). [9] A. Dimoulas, D. P. Brunco, S. Ferrari, J.W. Seo, Y. Panayiotatos, A. Sotiropoulos, T. Conard, M. Caymax, S. Spiga, M. Fanciulli, Ch. Dieker, E. K. Evangelou, S. Galata, M. Houssa, M. M. Heyns, Thin Solid Films (2007). [10] A. Ritenour, A. Khakifirooz, A. D. Antoniadis, R. Z. Lei, W. Tsai, A. Dimoulas, G. Mavrou, Y. Panayiotatos, Appl. Phys. Lett. 88, (2006). [11] N. Lu, W. Bai, A. Ramirez, C. Mouli, A. Rienour, M. L. Lee, D. Antoniadis, D. L. Kwong, Appl. Phys. Lett. 87, (2005). [12] J. W. Seo, C. Dieker, J.-P. Locquet, G. Mavrou, A. Dimoulas, Appl. Phys. Lett. 87, (2005). [14] H. Kim, P. C. McInyre, Chi O. Chui, K. C. Saraswat, M.-H. Cho, Appl. Phys. Lett. 85, 2902 (2004). [15] G. Tallarida, S. Spiga, C. Wiemer, M. Fanciulli, S. Kremmer, C. Teichert, in preparation. [16] C. Wiemer, S. Ferrari, M. Fanciulli, G. Pavia, L. Luterotti, Thin Solid Films 365, 82 (2000). [17] H. Kim, C. O. Chui, K. C. Saraswat, P. C. McIntyre, Appl. Phys. Lett. 83, 2647 (2003). [18] M. Perego, G. Scarel, M. Fanciulli, I. L. Fedushkin, A. A. Skatova, Appl. Phys. Lett. 90 (2007). 2.6 Epitaxial HfO 2 on high-mobility semiconductors: theory and experiment C. Wiemer, A. Debernardi, G. Scarel, M. Perego, M. Fanciulli CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy Despite the large lattice parameter mismatch, we observed a preferential orientation of HfO 2 films grown by atomic layer deposition (ALD) on Ge(001) substrate at 375 C using HfCl 4 and H 2 O as precursors [1]. The Si lattice parameter (5.43 Å) is more favourable than the one of Ge (5.66 Å) to promote epitaxial growth of HfO 2. Nevertheless, no epitaxial growth of HfO 2 films on Si is reported. The successful occurrence of local epitaxial growth of HfO 2 films on 89

91 Ge might instead be related to the absence of any amorphous interfacial oxide layer between the high-κ and the semiconductor. The lattice parameter of the substrate affects then the orientation of the HfO 2 grains. Films in the 6-21 nm thickness range grown using HfCl 4 and H 2 O on Ge (001) at 375 C exhibit the majority of the monoclinic {001} planes parallel to the Ge (001) surface [2]. The in-plane orientation of the monoclinic (001) planes is explored by a complete phi scan around the (013) asymmetric reflection (Figure 1). The elongated spots at phi = 0, 90, 180 and 270 are due to the Ge{311} reflections, whereas the four sharp spots at 45, 135, 225 and 315 are due to the monoclinic (013) and (103) HfO 2 planes. This finding suggests that there are two preferential in-plane orientations. Since the relative intensities of the four sharp spots are equal for all phi angles, we conclude that the growth of HfO 2 monoclinic (001) crystals oriented parallel to the Ge(001) plane takes place with the same probability in the two cube on cube in-plane orientations: Ge[100]//HfO 2 (100) and Ge[100]//HfO 2 (010). ones, (the estimated tensile stress is about -170 Kbar), according to the relaxation mechanism predicted by theory. We calculated that this new epitaxial phase -that has no bulk counterpart - has a dielectric constant equal to 11 along the growth direction, i.e. it is comparable to the bulk monoclinic one. According to our picture, thin epitaxial HfO 2 films will allow to control at the atomic level their interface with the semiconductor, as it is required for the integration of high-κ dielectrics in future complementary-metal-oxide-semiconductor devices. Figure 2. Structural and electronic (yellow isosurface sorrounding oxygen atoms) configuration of epitaxial HfO 2. Figure 1. Phi scan around the 013 reflection of monoclinic HfO 2 obtained on the 15 nm thick film grown on Ge(001). By state-of-the-art first principle calculations within the plane-wave pseudopotential method, we identify the microscopic mechanism responsible for the oriented growth of HfO 2 on Ge (001) surface. The observed preferential orientation of the monoclinic structure is related to the relaxation of an epitaxial HfO 2 phase, different from the monoclinic one, when a critical thickness is reached [3] (Figure 2). In fact, as revealed by X-ray diffraction analysis, the preferentially oriented monoclinic structure has the in-plane axis presenting the same orientation as the in-plane axis of the epitaxial phase. Further, the inplane lattice parameters of the (001) oriented grains are considerably larger than the bulk baddeleyite Due to the similarity between the lattice parameter of Ge and GaAs (5.65 Å), a preferential orientation of monoclinic HfO 2 crystals might be expected also on GaAs (001) clean surfaces. Actually, local epitaxy on GaAs (001) is also observed in HfO 2 films grown by ALD using the same precursor combination as the above mentioned one. The GaAs surface is cleaned using two different methods, one consisting in a 30 s immersion in a dilute HF solution, and the other consisting in a 2 min immersion in a HCl based solution, which is reported to be very effective in removing the native oxide. Combined X-ray photoelectron spectroscopy and X-ray reflectivity analysis show that HCl cleaning provides an As-richer surface, with less As oxides and a minor Ga 2, less significant than the one revealed on the GaAs surface cleaned in the HF solution. However, the HCl-based cleaning leaves a Cl residue at the interface and a ~0.5 Å rougher GaAs surface. Different deposition temperatures of the HfO 2 have been explored, ranging between 300 C and 425 C. For all these deposition temperatures, the GaAs/HfO 2 interface is rougher when the surface is prepared by HCl than when it is prepared by HF-dip. An amorphous interlayer is only present when the growth takes place at 300 C on the HF-dip-prepared GaAs surface, as evidenced by transmission electron 90

92 microscopy and electronic density profiles obtained from the fitting of XRR data [4] (Figure 3). On the other hand, for films grown at 375 C, the absence of an interlayer is supported by XPS measurements that do not reveal any Ga or As oxides between the HfO 2 layer and the GaAs substrate [3]. Figure 3. Electronic density profile of films grown at 300 C on GaAs with two surface preparations: HF-dip (green), HCl solution (red). Films grown at 300 C are mostly amorphous on GaAs surfaces prepared both using HF and HCl. Films grown at higher temperatures are crystallized. For the same thickness of the HfO 2 layer, the surface preparation does not influence the preferential orientation of the HfO 2 crystallites. On the other hand, a little increase of the (001) preferential orientation is found by increasing the growth temperature from 375 C to 425 C, as reported in Figure 4, where X-ray diffraction data taken in Bragg -Brentano configuration are shown. Actually, the desorbtion of mixed Ga and As oxides that might have formed during the ALD cycles should be enhanced at higher deposition temperatures. Figure 4. X-ray diffraction data of 9 nm thick HfO 2 films grown on a HF-dip prepared GaAs surface at 375 C (blue) and 425 (magenta). Figure 5 compares the X-ray diffraction Bragg- Brentano spectra of 18 nm thick HfO 2 grown on Ge and on GaAs at 375 C. The same epitaxial relationships observed for HfO 2 films on Ge hold for those grown on GaAs. However, due either to some Ga oxides formed on the GaAs surface and/or to the different oxidation states of Hf (4) on one hand, and Ga (3) and As (3) on the other, the monoclinic {001} out of plane orientations of the HfO 2 films are less developed on GaAs than on Ge. In particular, the monoclinic (001) reflection is much less intense for HfO 2 films grown on GaAs than on Ge, further suggesting that the effect of the substrate lattice parameter is screened by some effect related to the chemistry at the interface. Figure 5. X-ray diffraction data of 18 nm thick HfO 2 films grown on a HF-dip prepared GaAs (green) and Ge (red) surfaces. [1] S. Spiga, C. Wiemer, G. Scarel, G. Seguini, M. Fanciulli, A. Zenkevich, and Yu. Lebedinskii, in Advanced gate stacks on high mobility semiconductors, Eds. A. Dimoulas, E. Gusev, P. C. McIntyre, and M. M. Heyns, Series on Materials Science - Springer-Verlag (2007). [2] S. Spiga, C. Wiemer, G. Scarel, G. Tallarida, G. Scarel, S. Ferrari, G. Seguini and M. Fanciulli, Appl. Phys. Lett (2005). [3] A. Debernardi, C. Wiemer, M. Fanciulli, submitted. [4] C. Wiemer, G. Scarel, M. Fanciulli, Nice (France) EMRS May 29 th -June2 nd 2006, oral presentation. 91

93 2.7 Epitaxial Gd 2 films on Ge A. Molle 1, M. N. K. Bhuiyan 1, G. Tallarida 1, C. Wiemer 1, M. Fanciulli 1, G. Pavia 2 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy 12 STMicroelectronics, Agrate Brianza, Italy The continuous efforts in developing novel materials with high dielectric constant (high-κ) on non-standard substrates addresses the concomitant demands of higher mobility channels and of an ultimate gate oxide down-scaling for advanced electronic devices. With these issues in mind, in the present work we report on the epitaxial growth of Gd 2 thin films on Ge(001) substrates by focusing on the structural details of the oxide film [1] and of the interface with the Ge substrate [2]. The Gd 2 was grown in the deposition stage of a multi-chamber ultra high vacuum (UHV) system (base pressure: 1x10-10 mbar) equipped with a Reflection High Energy (30 kev) Electron Diffraction (RHEED) apparatus. The system also includes an analysis chamber for the in situ X-ray Photoelectron Spectroscopy (XPS) provided by a nonmonochromatic Mg Kα radiation source ( ev). The Ge substrates were prepared by several cycles of Ar ion sputtering at room temperature and annealing at 500 C then obtaining a high degree of structural order and chemical purity on the Ge(001) surface [3]. The Ge initial condition is illustrated in Figure 1(a) by a streaky RHEED pattern endowed with remarkable half-order streaks, (-½,0) Ge and (½,0) Ge, indicating a well-defined (2x1) reconstruction on the flat Ge(001) surface. The Gd 2 films were deposited via a pure Gd beam with a rate of 0.10 nm/min (pre-calibrated by a quartz balance thickness monitor) at the growth temperature T g =450 C with an O 2 partial pressure p[o 2 ]= mbar. The structure of the Gd 2 surface was monitored by real time RHEED along the [110] Ge surface direction. In particular, Figure 1(b) illustrates the in situ RHEED patterns of a Gd 2 film with thickness t ox =2.7 nm. The symmetry of the initial (2x1) Ge(001) surface disappears after few seconds of deposition and twofold spots are observed in the RHEED pattern in Figure 1(b) around the central component which demonstrates the emergence of a cubic (110) registry with two orthogonal domains (bixbyite structure) in the Gd 2 film even at the sub-monolayer growth regime consistently with the Gd 2 /Si(001) growth [4]. To support this evidence, the ex situ XRD scan in Figure 2 ensures that the film grows epitaxially cubic (110) in agreement with the RHEED in Figure 1(b). As shown in Figure 2, only the cubic (440) plane is identified in the Bragg-Brentano line, while complete phi scans around the [111] direction reveal the presence of the two different, equally probable, in-plane domain rotations, in agreement with what reported for the growth of Gd 2 on Si [4] and on GaAs [5]. From the d hkl of the (110) and the (111) planes, the in-plane lattice parameter nm is deduced, whereas the out-of-plane lattice parameter is nm. In order to accommodate on the Ge(001) surface, the cubic (110) Gd 2 structure turns out to be tetragonally distorted with the in-plane lattice parameter closely matching that of the substrate (only 0.5 % smaller) and the out-of-plane lattice parameter being only 1.7% lower than that of the relaxed Gd 2 bulk [6]. The chemical and compositional details of the Gd 2 film are investigated by means of the Gd 4d XPS line for a film with thickness of t ox =16.5 nm as shown in Figure 3(a). The Gd 4d line of a pure 4 nm-thick Gd film is also reported as a comparison in order to (a) (b) Figure 1. RHEED pattern of the as-prepared (2x1) Ge(001) surface (a) and of a Gd 2 film with thickness t ox =2.7 nm (b). In the latter pattern the two orthogonal domains of the bixbyite structure are identified by the two reciprocal surface lattice parameters, g //,1 and g //,2. 92

94 intensity (arb. units) (a) Gd 4 d Gd 2 Gd 9 D D (b) Ge 3 d x 4 pure Ge Gd 2 /Ge x 4 Gd/Ge Figure 2. XRD analysis in Bragg Brentano configuration of 2.7 nm-thick sample. Inset: complete maps around the two [111] ox directions, corresponding to the two differently oriented domains. understand how the Gd-O bonding in the Gd 2 film affects the Gd 4d XPS shape profile. In agreement with ref. [7], the spectrum of the Gd 2 presents a shift of the 9 D and the 7 D components towards higher binding energies (BEs) and an overall broadening of the widths. In addition, the deep minimum between the two components, typical of the metallic character of the pure Gd film, smoothens in the Gd 2 case. By comparing the XPS integrated intensities of the Gd 4d and of the O 1s lines the oxide composition turns out to be fully compatible with the Gd 2 stoichiometry. To explore the Ge bonding configuration at the interfacial level, an in situ XPS Ge 3d line analysis was performed on the 2.7 nm-thick Gd 2 film. In order to check the presence of Ge-O or Ge-Gd bondings, the XPS Ge 3d line of the Gd 2 film in Figure 3(b) is compared to the Ge 3d line of an asprepared Ge(001) surface and of a 4 nm-thick Gd film grown on Ge(001). The XPS spectrum of the asprepared Ge surface presents an intense Ge 3d peak at binding energy BE=29.2 ev. A clear shift of the Ge 3d line towards lower BE s is observed in case of the 4 nm -thick Gd film on Ge(001) due to the Gd-Ge bonding. The XPS Ge 3d line registered on the Gd 2 film has the same peak position as that obtained on the as-prepared Ge surface (Figure 3(b)). No additional components due to the Ge oxidation states are observed so that the formation of a Ge-O as well as Ge-Gd bonding can be here ruled out in the Gd 2 film. Overall, the XPS analysis demonstrates that no IL forms in the Gd 2 growth on Ge(001) as observed for the Gd 2 on Si(001) [4] and on GaAs(001) [5]. Time of flight secondary ion mass spectroscopy measurements support this picture, indicating sharp interface between the Gd 2 film and the Ge substrate. This is consistent with the high resolution (c) nm Binding Energy (ev) Gd2O3 Ge Figure 3. (a) XPS Gd 4d line of a 16.5 nm-thick Gd 2 film (circles). As a comparison, the XPS Gd 4d line of a 4 nm-thick pure Gd film grown on Ge(001) is also reported (grey dots); (b) Comparison among the XPS Ge 3d line of an as-prepared Ge(001) surface, a 2.7 nm-thick Gd 2 film on Ge(001) and a pure Gd film on Ge(001) (the first and the third XPS scans have been divided by a factor 4); (c) Cross sectional TEM image of the interface region of a Gd 2 film grown on Ge(001). The cut was done along the [110] plane of the Ge substrate. TEM image of the Gd 2 /Ge interface region in Figure 3(c), which emphasizes how the crystalline registry of the film superimposes on the bulk registry without forming any interface layer. Indeed, the strong contrast between the Ge and the oxide region evidences an atomically abrupt interface. Upon increasing the thickness, the oxide undergoes a phase transition from the cubic (110) to a mainly monoclinic (100) structure. The identification of the major monoclinic (100) phase is supported by XRD analysis, which identifies the largely dominant monoclinic (600) Bragg peak over the cubic (440) component [1]. This work has been partially supported by the European Project ET4US. [1] A. Molle, C. Wiemer, M. N. K. Bhuiyan, G. Tallarida, G. Pavia, and M. Fanciulli, in press on Appl. Phys. Lett. [2] A. Molle, C. Wiemer, M. N. K. Bhuiyan, G. Tallarida, G. Pavia and M. Fanciulli, Appl. Phys. Lett (2007). [3] A. Molle, M. N. K. Bhuiyan, G. Tallarida, and M. Fanciulli, Appl. Phys. Lett. 89, (2006). [4] A. Laha, H. J. Osten, A. Fissel, Appl. Phys. Lett. 89, (2006). [5] M. Hong, J. Kwo, A. R. Kortan, J. P. Mannaerts, A. M. Sergent, Science 283, 1897 (1999). [6] Inorganic Crystal Structure Database, files 96207, Fachinformationzentrum Karlsruhe (2005). [7] P. Terzieff and K. Lee, J. Appl. Phys. 50, 3565 (1979). 93

95 3. Innovative Technologies in Non-Volatile Memories 3.1 Si nanocrystals in a dielectric matrix: synthesis and characterization 3.2 Metallic nanocrystals embedded in SiO 2 : low temperature CEMS characterization 3.3 Structural and functional properties of phase change materials grown by sputtering and MOCVD 3.4 Thermal characterization on phase change materials 3.5 Binary oxides for resistive switching non volatile memories 3.6 Electrical and chemical investigations of ZnO layers grown by Atomic Layer Deposition for selectors in crossbar non volatile memories The Non Volatile Memories (NVM) became one of the major markets in the microelectronic field, due to the growing popularity of memory hungry consumer electronics devices for multimedia and communication technologies. At the moment Non Volatile Memories (NVM) relies almost exclusively on floating gate technology. Even if they still present a lower density, in terms of Megabytes/cm 2, a higher cost, in terms of pure cost/bit, when compared to other mass-storage media (magnetic and optical), they present also some critical advantages, like the lack of any mechanical parts, which results in stronger ruggedness, lighter weight, smaller formfactor, better reliability and lower power dissipation, which make them ideally suited for portable applications. Even if Intel expressed its confidence in the possibility to extend floating gate technology to the 45 and possibly 32 nm nodes, physical limitations exist that could further reduce the margins for cell size reduction, unless substantial progress is made in critical areas. These limits are essentially related to the difficulty in reducing tunnel and interpoly dielectric thickness, due to trap-assisted leakage currents. At the MDM Laboratory we are exploring new NVM concepts and architectures that might eventually overcome the intrinsic limitation imposed by the floating gate technology. 94

96 Most of the MDM activities in this field were carried out in the frame of several European Projects spanning from FP5 to FP6. An evolution of the floating gate concept is represented by a memory type that replaces the floating gate based on the polysilicon sandwitched between two oxides layer, with an array of metallic or semiconducting nanocrystals embedded directly in the gate dielectric. We explored in particular the use of nanocrystals of Si and Sn embedded in a dielectric matrix obtained by Ion Beam Synthesis (IBS), i.e. ion implantation followed by thermal treatment. The focus of the paragraph 3.1 is on the fabrication and chemical and structural characterization of Si nanocrystals in SiO 2 and high-κ oxides. The material science issues related to the synthesis by IBS of the Si nanocrystals as well as their integration in real memory devices were the subject of the European Project NEON. The paragraph 3.2 focuses on the characterization of Sn nanocrystals in SiO 2 by conversion electron Mössbauer spectroscopy (CEMS). An emerging NVM concept not based on the floating gate technology is represented by the Phase Change Memory (PCM). The PCM cell is essentially a resistor of a thin-film chalcogenide material (Ge 2 Sb 2 Te 5 -GST) with a low-field resistance that changes by orders of magnitude, depending on the phase state of the GST in the active region (i.e., crystalline or amorphous). The development of new phase change materials and their integration in PCM devices is the subject of the European Project CHEMAPH The paragraph 3.3 deals with structural and functional properties of phase change materials grown by sputtering and MOCVD, while the paragraph 3.4 relates to the thermal characterization on phase change materials. Another very interesting NVM technology is constituted by the so called resistive switching memories. This is a very broad class of memories that includes all the devices that store the data as a variation of the resistance read between the top and bottom electrode. There are two key elements that characterize such type of memory: the bit storage element that determines the bit state that is usually constituted by the resistive switch material and the selection element that can be either be constituted by a transistor or by a diode and allows to uniquely address a specific bit in a memory array. In paragraph 3.5 we describe the use of binary oxides for resistive switching in NVM, that is the subject of the European Project EMMA, while in paragraph 3.6 we describe the realization of the selection elements based on diodes for NVM using ZnO that is the subject of the European Project VERSATILE. 95

97 3.1 Si nanocrystals in a dielectric matrix: synthesis and characterization M. Perego 1, S. Spiga 1, M. Fanciulli 1, C. Bonafos 2, and G. Benassayag 2 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy 12 MAT Group, CEMES-CNRS, Toulouse, France A decrease of the injection distance down to 5 nm is obtained in 10 nm thick SiO 2 films by increasing the implantation energy up to 3 kev. Smaller values (2 nm) of the injection distance can be achieved by implanting 1keV Si ions in thinner oxides [3]. The combination of experimental and theroretical studies allows to explain these experimental results [3]. Transferring this methodology to high-κ material, like Al 2 and HfO 2, different problems have been identified [6]. First in the 1 kev implanted samples a very thick amorphous silicon region has been observed in the substrate (Figure 1). The possibility to use semiconducting or metallic nanocrystals embedded in a SiO 2 matrix as charge storage elements in novel non volatile memory devices has been widely explored in the last ten years. The replacement of the continuous polysilicon layer of a conventional flash memory device by a 2-dimensional Si nanoparticle array presents several advantages but the fundamental trade-off between programming and data retention characteristics has not been overcome yet [1,2]. The main problem is the limited retention time basically due to charge loss by leakage current through the ultrathin SiO 2 tunneling dielectric. A longer retention time can be achieved by increasing the tunnel oxide thickness. This however implies higher operating voltages and consequently a reduced write/erase speed. Using high-κ materials for tunnel and/or gate oxide in principle it is possible to achieve the goal of a low voltage non volatile memory device. The high dielectric constant of these materials allows to use thicker tunnel oxide reducing leakage current. Several approaches have been explored to synthesise ordered nanocrystal arrays in SiO 2 but the transfer of these methodologies to the synthesis of 2-d array of nanocrystals in high-κ materials is not trivial. Among the different high-κ materials under investigation, Al 2, HfO 2 and their alloys are probably the most widely studied and are also considered as very promising candidates for the integration in ultra-scaled commercial devices. Thus, we focus on the possibility to synthesize silicon nanocrystals in Al 2 and HfO 2 using two different techniques that have already been widely tested for the synthesis of Si nanocrystals in a SiO 2 matrix; ultra low energy ion beam synthesis (ULE- IBS) and e-beam evaporation of multilayer structures. ULE-IBS has been demonstrated to be a very powerful technique, that allows to place the nanocrystal array at a controlled distance from the SiO 2 /Si interface [3-5]. Figure 1. HRTEM images of a 1 kev Si implanted (a) and implanted and 900 C annealed Al 2 sample (b). In the as implanted sample a 7.6 nm thick glassy region consisting of amorphous silicon can be observed. Re-crystallization can be observed after annealing. Extended defects are present below the re-crystallized silicon. This amorphous layer re-crystallizes with bulk silicon parameters after annealing, but extended defects have been observed below the re-crystallized silicon. Second, implanted silicon has been either massively oxidized or incorporated within the high-κ matrix during the high temperature thermal treatment. According to different experimental evidences, a moisture absorption process in the heavily damaged near surface region of the dielectric film was suggested to explain this phenomenon. Finally a Si diffusion process in the dielectric matrix has been observed at high temperatures for the Si implanted Al 2 samples. Thermal evaporation and annealing of SiO x /SiO 2 superlattices is an alternative approach explored to form Si ncs multilayers for optoelectronics application. This methodology is intriguing also for memory devices because ordered arrays of nanocrystals, with high density and limited dispersion in size, can be precisely positioned in a dielectric matrix. We studied the formation of ordered arrays of Si ncs using e-beam evaporation of SiO/SiO 2 multilayer and subsequent annealing at high temperatures ( C) [7]. The formation of Si nanocrystals layers has been confirmed by high resolution transmission electron microscopy 96

98 Figure 2. HREM image showing a clear evidence of the presence of crystalline clusters embedded in a silicon oxide matrix. The two bands have been obtained by deposition of a SiO 2 /SiO multilayer structure followed by a 60-min thermal treatment at 1000 C in N 2. The images, with well resolved lattice planes, of these Si ncs are shown in the insets. (Figure 2). Capacitance-voltage measurements show charge trapping phenomena related to the presence of the Si ncs in the oxide matrix (Figure 3). Unfortunately the e-beam evaporation and subsequent annealing of multilayer structures with a SiO thin film embedded in a high-κ matrix was not effective in synthesizing the Si nanocrystals [6]. Si amorphous clusters in a Al 2 matrix were obtained after a 900 C thermal treatment (Figure 4). Higher temperatures were considered in order to induce the crystallization of the silicon clusters but a silicon out-diffusion mechanism was activated during the thermal treatment preventing nucleation and growth of the nanoclusters. No clusters at all were obtained in HfO 2 /SiO/HfO 2 multilayer. This indicates that a partial reoxidation of the SiO film occurs during annealing. This parasitic oxidation has been related to an excess of oxygen in the HfO 2 films. The use of a SiO 2 capping layer on top of the HfO 2 stack has been observed to prevent this phenomenon. In conclusion the synthesis of Si nanocrystals in Al 2 and HfO 2 matrix by ULE-IBS or e-beam evaporation and subsequent thermal treatment are prevented by two main factors. The first problem is related to the fact that at high temperatures Si diffuses in the oxide matrix. Figure 3. High-frequency capacitance-voltage curves (acquired at 500 khz) for the two different MIS structures. Structure including Si nanocrystals evidences a large clockwise hysteresis which might be related to the charging effects of Si nanoclusters embedded in the SiO 2. Figure 4. ToF-SIMS depth profiles of the as-deposited (c) and of the 900 C (d) and 1000 C (e) annealed Al 2 /SiO/Al 2 multilayer structure. The higher intensity of the Si 2 - signals after the 900 C thermal treatment suggests the formation of Si nanoclusters. Defocused bright field TEM image (a) confirms the presence of a Si rich region in the 900 C annealed sample. The HRTEM image (b) does not show any crystalline structure in this silicon rich layer. 97

99 The second problem is related to parasitic oxidation of Si during the thermal treatment due to the excess of oxygen present in the high-κ matrix. [1] P. Normand, E. Kapetanakis, P. Dimitrakis, D. Skarlatos, K. Beltsios, D. Tsoukalas, C. Bonafos, G. Ben Assayag, N. Cherkashin, A. Claverie, J. A.Van Den Berg, V. Soncini, A. Agarwal, M. Ameen, M. Perego, M. Fanciulli, Nuclear Instr. Meth. B 216, (2004). [2] P. Dimitrakis, E. Kapetanakis, D. Tsoukalas, D. Skarlatos, C. Bonafos, G. Ben Assayag, A. Claverie, M. Perego, M. Fanciulli, V. Soncini, R. Sotgiu, A. Agarwal, M. Ameen, C. Sohl, P. Normand, Solid-State Electronics 48, (2004). [3] C. Bonafos, M. Carrada, N. Cherkashin, H. Coffin, D. Chassaing, G. Ben Assayag, A. Claverie, T. Muller, K. H. Heinig, M. Perego, M. Fanciulli, P. Normand, D. Tsoukalas, J. Appl. Phys. 95, (2004). [4] N. Cherkashin, C. Bonafos, H. Coffin, M. Carrada, S. Schamm, G. Ben Assayag, D. Chassaing, P. Dimitrakis, P. Normand, M. Perego, M. Fanciulli, T. Muller, K. H. Heinig, A. Claverie, Physica Status Solidi (c) 2(6), (2005). [5] M. Perego, S. Ferrari, M. Fanciulli, G. Ben Assayag, C. Bonafos, M. Carrada, A. Claverie, J. Appl. Phys. 95, (2004). [6] M. Fanciulli, M. Perego, C. Bonafos, A. Mouti, S. Schamm, G. Benassayag, Advances in Science and Technology 51, (2006). [7] M. Perego, M. Fanciulli, S. Spiga., C. Bonafos, N. Cherkashin, Materials Sci. and Engineering C, 26, (2006). 3.2 Metallic nanocrystals embedded in SiO 2 : low temperature CEMS characterization could be very different when compared to their bulk counterpart, due to the large surface-to-volume atoms ratio. For instance, the melting properties of Sn metallic isolated nanoparticles are size dependent [4]. Thin SiO 2 films (85 nm thick) are implanted with 119 Sn + ions at room temperature (RT), with energy of 80 kev and fluence 1x10 16 cm -2. Rapid thermal annealing follows ion implantation, in order to both recover the implantation damage, and obtain a thermally activated and controlled redistribution of the implanted ions by phase separation [5]. Here, we consider a set of samples annealed at 900 C for 30 s, 900 C for 120 s, and 1100 C for 30 s. A detailed study on the local structure of Sn-implanted in thin SiO 2 films, has been performed by means of conversion electron Mössbauer spectroscopy (CEMS) at room temperature (RT), extended X-ray absorption spectroscopy (EXAFS), and transmission electron microscopy (TEM) [5]. In the as-implanted samples all the Sn atoms are oxidized (with both Sn 2+ and Sn 4+ oxidation states present), while annealing induces the formation of metallic β-sn nanocrystals, leaving a fraction of Sn atoms oxidized. Figure 1 shows a TEM image of the sample annealed at 900 C for 30 s. It is evident the presence of clusters of different diameters, in the range 2-10 nm. High-resolution TEM (not shown) evidences that the biggest clusters are crystalline and associated with β-sn. Small amorphous cluster (2-3 nm) are also shown by TEM either in the as-implanted and annealed samples. R. Mantovan, S. Spiga, A. Debernardi, and M. Fanciulli CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy Metallic and semiconducting nanocrystals embedded in dielectric materials are attracting great interest due to their novel electrical and optical properties. In particular, for Sn nanocrystals formed in SiO 2 by ion implantation, single electron charging effects [1] and new luminescence properties [2] have been observed. Memory devices using nanocrystals as charge storage elements have been also demonstrated [3]. The physical properties of low-dimensional systems Figure 1. TEM image of the sample annealed at 900 C, 30 s. We study the local atomic structure of the Sn atoms by performing CEMS at different temperatures, i.e. 295 K and 120 K (LT). CEMS performed at different temperatures allows the determination of the Debye- 98

100 Waller factors f and the Debye temperatures (Θ D ) of the Sn-related phases, making their identification more reliable. In order to perform LT-CEMS, we have developed an experimental set-up based on a sealed parallel-plate avalanche counter (PPAC), which allows extending CEMS down to 120 K [6]. This novel PPAC has been tested by measuring Θ D of a reference β-sn bulk sample, having a thin native SnO 2 oxide on its surface, and the results are 121 K and 215 K respectively [6]. Figure 2 shows the RT and LT spectra for the sample annealed at 900 C for 30 s. The isomer shifts are given relative to CaSn at RT. CEM-spectra are fitted in terms of four contributions, related to the oxidized Sn 2+ (a), Sn 2+ (b), Sn 4+ phases, and to the metallic β-sn phase [5,7]. Relative emission 900 C 30 s 5% 20% RT LT Sn 2+ (b) 2+ Sn Velocity (mm/s) 4+ Sn (a) Figure 2. CEM-spectra at RT and LT for the sample annealed at 900 C for 30 s. The spectra are fitted using four components: Sn 2+ (a), Sn 2+ (b), Sn 4+, and β-sn. by normalizing the spectral fractions to the obtained f factors, Eq. (1), for all the Sn-related phases. In contrast to the results based on CEMS at RT [5] for the sample annealed at 900 C for 30 s, it is found that the Sn 2+ (a) phase is predominant (61%) with respect to the Sn 2+ (b) one (30%) [7]. Less than 20% of Sn atoms are in the Sn 4+ local electronic configuration. We relate the Sn 4+ oxidation state to oxidized Sn atoms close to the SiO 2 surface [7]. The two different Sn 2+ local configurations have distinct Θ D values. Θ D for Sn 2+ (a) is 137(7) K, well below the values reported for crystalline and amorphous SnO, which are 203(1) K and 181(2) respectively [8]. On the other hand, for Sn 2+ (b) the calculated Θ D is 168(13) K, being very close to that of the amorphous SnO compound. For Sn 2+ (a) and Sn 2+ (b), the hyperfine parameters are different than those characterizing SnO [8]. We suggest that the Sn 2+ (a) contribution in the CEM-spectra is related to the small (2-3 nm) amorphous SnO clusters, also detected by TEM (Figure 1), in which the Sn atoms are located in a highly disordered local environment [7]. The component b might be related to Sn atoms dispersed in the SiO 2 matrix. After the implantation and annealing processes, the Sn atoms might be present into the SiO 2 matrix in substitutional position. The Sn local environment is similar to that one of the SnO amorphous compound, as confirmed by the measured Θ D and hyperfine parameters [7]. Another possible explanation is that the two distinct components a and b are due to different SnO x mean cluster size. Following annealing processes at higher temperatures or for longer time, CEMS revealed a decrease of the spectral intensity related to component Sn 2+ (b) [5]. This is possibly due to the thermal instability of SnO, which may dissociate forming β-sn clusters. From Figure 2, the variation of the resonant absorption area A with varying the temperature is well visible. The Θ D is calculated according to the expression of f in the Debye model, Eq. (1). ΘD 2 T 3E R T x f = exp Θ Θ dx x 2k D D 0 e 1 where E R is the recoil energy of the γ-ray, k is the Boltzman constant, and T is the samples temperature. Since, for each component, f is proportional to A, we determine the Θ D for the Sn-related phases from the temperature dependence of A. We use the measured Θ D to make a comparison between the different phases. The real site-populations can be calculated 1) The mean β-sn-cluster diameters d, as determined by TEM, are 7.4(6), 10.9(9), and 16.9(7) nm for the samples annealed at 900 C for 30 s, 900 C for 120 s, and 1100 C for 30 s respectively [5]. By using a simple model, which is based on CEMS results at RT and Rutherford backscattering spectroscopy data about the total amount of Sn atoms in the samples, we estimate f for the β-sn nanocrystals having different size [5]. We find a considerable reduction of f for the smallest clusters, while for the biggest ones it approaches the value reported for the bulk β-sn compound [5]. We study the size dependence of f in β-sn nanocrystals, by performing CEMS at different temperatures [9]. The Debye-Waller factor is intimately related to the dynamical processes involving the nuclei during the γ-ray emission. We find that f in β-sn nanocrystals embedded in SiO 2 matrix is strongly size dependent. 99

101 Debye-Waller factor f fbulk (CEMS) RT f CEMS f Ref. [5] f theory F. Boscherini, F. d Acapito, B. Schmidt, R. Gröetzschel, and A. Mücklich, Phys. Rev. B 68, (2003). [6] R. Mantovan and M. Fanciulli, Review of Scientific Instruments, in press. [7] R. Mantovan, S.Spiga, and M. Fanciulli, Hyperfine Interactions 165, 69 (2005). [8] G. S. Collins, T. Kachnowski, N. Benczer-Koller, and M. Pasternak, Phys. Rev. B 19, 1369 (1979). [9] R. Mantovan, A. Debernardi, and M. Fanciulli, Phys. Rev. B, in press Figure 3. Size dependence of the Debye-Waller factor in β-sn nanocrystals at RT=295 K. (Blue) stars: experimental data, as obtained by fitting the temperature variation of A with the Debye model; (green) triangles: taken from [5]; (red) circles: theoretical results. The (red) continuous line is reported as a guide for the eyes. The horizontal dashed lines indicate the f values for bulk β-sn, as measured with CEMS. The largest nanocrystals have f approaching the bulk value, while (at RT) the smallest clusters display a 60% decrease, which corresponds to a 30% increase of the mean square displacement. The experimental results are compared with calculation based on a continuous model in which the relevant parameters are determined ab initio [9]. This model considers spherical nanocrystals, and is based on the elastic body approximation including surface relaxation effects. Figure 3 shows the comparison between the CEMS results and the theoretical calculation [9]. We find accordance between the Mössbauer experiment and the theoretical results. The size dependence of f is due to the remarkable contribution of the surface vibrational modes in determining the vibrational density of states (VDOS) in the smallest nanocrystals, i.e. a surface relaxation effect. The size dependence of f is more pronounced as the surface-to-volume atoms ratio becomes higher. For large particles, i.e. d 16 nm, the nanocrystals VDOS approaches the bulk one. [1] A. Nakajima, T. Futatsagi, H. Nakao, T. Usoki, N. Horiguchi, and N. Yokoyama, J. Appl. Phys. 84, 1316 (1998). [2] J. M. Lopes, F. C. Zawislak, P. F. P. Fichter, F. C. Lovey, and A. M. Condò, Appl. Phys. Lett. 86, (2005). [3] T. Futatsagi, A. Nakajima, and H. Nakao, FUJITSU Sci. Tech. J. 34, 142 (1998). [4] T. Bachels, H. J. Güntherodt, and R. Schäfer, Phys. Rev. Lett. 85, 1250 (2000). [5] S. Spiga, R. Mantovan, M. Fanciulli, N. Ferretti, 3.3 Structural and functional properties of chalcogenide materials grown by sputtering and MOCVD C. Wiemer 1, M. Fanciulli 1, C. Giessen 2, R. Bez 3, A. Pirovano 3, S. Rushworth 4, J. Siegel 5, C. N. Afonso 5, A. Abrutis 6, V. Plausinaitiene CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy Aixtron AG, Kackertstr, Aachen Germany STMicroelectronics, Agrate Brianza, Italy Epichem / SAFC HiTech Power Road, Bromborough, Wirral, CH62 3QF, United Kingdom Laser Processing Group, Instituto de Optica, CSIC, Serrano 121, E28006 Madrid, Spain Vilnius University, Faculty of Chemistry, Dep. of General and Inorganic Chemistry, Naugarduko 24, LT-2006 Vilnius, Lithuania Chalcogenide-based phase change materials, already applied as optical storage media, are nowadays evaluated to be implemented for electronic memory applications. They constitute the most promising candidate to scale non-volatile-memory technology beyond the Flash memory architecture. One of the technological issue that determines the integration of these new materials is the deposition method. Demonstrator chips based on phase change memories, with chalcogenide material grown by sputtering have already been proved. However, due to the continued down-scaling for nanoelectronic device architectures, greater control of film deposition over non-planar structures than achieved by sputtering is necessary. This would allow lower programming currents, which lead to improved performance and lower costs. We investigate the growth of chalcogenides materials by metal organic chemical 100

102 vapor deposition (MOCVD), a chemical-based technique that enables the production of thin films with better conformality, coverage, and stoichiometry control then those deposited by sputtering. In this contribution we compare structural and functional properties of sputtered-deposited Ge 2 Sb 2 Te 5 (GST) and chalcogenide materials grown by liquid-injection MOCVD. The structural properties of sputtered-deposited GST are addressed by X-ray reflectivity (XRR) and X-ray diffraction (XRD) analysis. Figure 1 shows the XRR data and simulation of a nominally 20 nm thick film on Si substrate. The film is so smooth that the thickness oscillation can be registered by XRR. The electronic density is close to the one reported for fully crystallized face centred cubic Ge 2 Sb 2 Te 5 (1.53 e - /Å 3 ). The roughness is below 2% of the thickness. A thin, low Figure 3. Temporal evolution of the reflectivity decrease of a Ge 2 Sb 2 Te 5 surface induced upon irradiation with a single laser pulse of 8 ns duration (amorphization, black) and reflectivity recovery by irradiation with consecutive laser pulses (re-crystallization, pink). density cap layer on top of the Ge 2 Sb 2 Te 5 is necessary to correctly simulate the XRR data. This layer can be related to some oxide developing on the surface of the chalcogenide. XRD analysis (Figure 2) confirms the crystallization in the face centred cubic phase, with lattice parameter extracted from the Rietveld refinement equal to 6.09 Å. However, not all the peak positions are correctly reproduced by the refinement, suggesting that some anisotropic strain might be present within the differently oriented grains. The average size of the diffracting domains, fixing the microstrain at 2.5% is 20.3 nm. The intensity of the diffracted peaks differs from the one reported for GST powder, suggesting a preferential orientation, or a different atomic arrangement and/ or stoichiometry. Figure 1. XRR data (red dots) and simulation of a nominal 200 nm thick GST layer grown by sputtering. Figure 2. XRD analysis of sputtered deposited GST: data (red) and simulation (black). The diffraction pattern as reported in ICSD database (blue) is also shown for comparison [1]. Up till now the most promising phase change material we achieved by liquid injection MOCVD is Sb doped with few Ge percent. The surface of the deposited layer is not smooth but shows a grainy microstructure. XRD analysis of this sample reveals only crystallization of the Sb crystals, an amorphous matrix is not excluded. Functional characterization of phase change materials is performed by real time reflectivity measurements of films grown on Si substrates covered with 100 nm SiO 2. Single pulse amorphization of sputtered GST was possible at high pulse energies for both ns and fs laser pulses. The inverse process, re-crystallization was only possible under multi-pulse irradiation and more pulses were required when using fs than ns laser pulses. Moreover, both transformation thresholds are considerably lower and the transformation time is shorter for fs than for ns laser pulses. Figure 3 101

103 shows the phase reversibility achieved by ns laser pulses. Starting from an as-deposited crystalline state, a single ns laser pulse is used to amorphize the exposed region, generating an amorphous spot. Re-crystallization is reached after irradiation with 4 consecutive laser pulses at, as it is clearly shown by the gradual increase of reflectivity. Irradiation using multiple ns laser pulses have been performed to investigate possible partial or complete phase transformation of Ge-doped Sb grown by MOCVD. Phase cycling was attempted by irradiating the same spot three times using different fluences. Several different fluences have been tried for all the pulses. The results shown in Figure 4 correspond to the fluences that lead to the best cycling performance in terms of optical contrast. This work is partially supported by the European Project Chemaph. [1] Inorganic Crystal Structure Database (ICSD), Fachinformationzentrum, Karlsruhe, 2006 fine n Thermal characterization of phase change materials J. L. Battaglia 1,2, A. Teren 1, C. Monguzzi 1, E. Varesi 3, R. Cecchini 3, C. Wiemer 1, R. Fallica 1, A. Kusiak 2, C. Rossignol 4, N. Chigarev 4, S. Cocco 1 Figure 4. White light microscopy images taken after 3 consecutive 8 ns laser pulses on the same region: left: after firs pulse, centre, after the second pulse and right, after the third pulse. While the phase transition during the 1 st pulse shows a complex reflectivity evolution, probably due to the initial grainy microstructure, the transitions during the following pulses are faster and simpler indicating the following mechanisms to occur. 2 nd pulse: fast decay (melting), fast increase (recalescence), slow decrease (solidification into the amorphous phase). 3 rd pulse: fast decay (melting), fast increase (crystallization). However, the initial fast reflectivity decrease which follows each pulse could also correspond to the evaporation of some material at the film surface. This evaporation would explain the much smoother surface observed at the centre of the irradiated areas (see above image after 3 rd pulse). This analysis evidences the need for optimising the composition of the phase change layer, in order to achieve optimum cycling conditions. In order to optimise the phase change procedure, the thickness of the material, as well as the nature of the buffer layer, should be carefully considered because they have a strong influence on the conditions of heat flow. Being related to thermal properties, these considerations apply to both optically and electrically induced phase transformations and are independent from the method used for depositing the phase change material CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy Laboratoire TREFLE, Ecole des Arts et Métiers, Talence Cedex, France STMicroelectronics, Agrate Brianza, Italy Laboratoire LMP, Université Bordeaux 1, Talence Cedex, France The knowledge of the thermal properties of Ge 2 Sb 2 Te 5 (GST) is required when simulating the heat transfer in a Non Volatile Memory device based on such a phase change material. Nevertheless, these properties are strongly dependent on the deposition process and it is of a major interest to know their dependence on the electrical, structural and optical properties. In order to obtain reliable values, thermal properties of the Ge 2 Sb 2 Te 5 are investigated using different ways: the 3ω technique, the photothermal radiometry (PTR) and the time domain thermoreflectance (TDTR). These methods are based on the same principle. The material, initially at stable and uniform temperature, is thermally excited from a heat flux and the transient temperature rise is measured, as a spatial averaged quantity, at the location of the heating. The advantage of the 3ω technique with respect to the photothermal methods (PTR and TDTR) is that it allows to measure with a good accuracy the heat flux and the average temperature of the heated area. Nevertheless, the 3ω technique requires a longer implementation time due to the narrow strip deposition, as well as to the realization of an electrical insulation between the strip and the GST in its crystalline phase. The 3ω and PTR 102

104 techniques mainly work at low frequency ( Hz for the 3w and Hz for the PTR) where the deposit (GST and additional layers) is viewed as a thermal resistance. The TDTR operates at very high frequencies ( Hz) where it is possible to see the heat diffusion in the deposit and thus to access to the thermal diffusivity or effusivity. To summarize, the 3ω and PTR techniques lead to measure the intrinsic thermal conductivity of the GST, whereas the TDTR permits measuring the thermal effusivity which is defined as the square root of the product between the heat capacity and thermal conductivity. The 3ω experimental technique is implemented at the MDM laboratory to measure the thermal conductivity of thin films prepared on a substrate. A gold narrow metal strip is deposited on the material to be characterized and a periodic driving current leads to Joule heating in the strip. The average temperature in the strip increases according to this heat source and heat diffuses in the SiO 2 /GST/substrate composite material (SiO 2 layer is required as an electrical insulator between the strip and the crystalline GST). The average temperature rise of the strip can be calculated from the measurement of the third harmonic of the voltage drop using a lock-in amplifier (see Figure 1). This measured temperature is used in order to identify the thermal resistance of the SiO 2 /GST stack from a model that describes the heat diffusion in the sample. A new heat transfer model has been developed [1] at MDM that takes into account the 2D geometrical configuration as well as the finite thickness of the silicon substrate for low-frequency behaviour. Amorphous and crystalline (face-centred-cubic and hexagonal) Ge 2 Sb 2 Te 5 of different thicknesses have been deposited using the DC magnetron sputtering technique on a silicon wafer. Each thickness has been carefully measured using SEM. This study leads to discriminate the intrinsic thermal conductivity of GST Figure 2. Measured (squares) and simulated (line) average temperature rise of the narrow strip for amorphous GST of different thicknesses. Figure 3. Experimental setup for the photothermal radiometry (PTR) method. Figure 1. Schematic description of the 3ω technique for thermal conductivity measurement of a thin film. within these three phases as well as the total thermal resistance at the interfaces. Figure 2 presents the results for amorphous GST. An accurate fit between the measured and simulated data, starting from the identified thermal conductivity, is obtained. The photothermal radiometry is used to measure the thermal conductivity of the amorphous and crystalline GST. This technique has been implemented at the TREFLE laboratory (ENSAM, France) and successfully applied for sub-micrometric layers [2,3]. The experimental setup, represented in Figure 3, is composed by two parts: the thermal excitation and the radiative heat flux measurement. The thermal excitation is generated on the surface of the sample 103

105 using a laser (UV and NIR laser diode according to the optical properties of the sample). The laser is modulated until 20 khz. Using an optical fibre, the laser beam is redirected perpendicular to the surface of the sample and focused by an objective. A nickel oxide (NiO) film, 50 nm thick, is deposited on the GST layer in order to increase the radiative absorption of the laser. Heat radiation from the heated area on the layer is measured using a HgCdTe based photovoltaic infrared detector (5-13 μm), cooled to liquid nitrogen temperature. A lock-in amplifier is used to measure the amplitude between the heat flux from the laser diode and the radiative heat with the IR detector, which is linearly proportional to the temperature on the heated area, as a function of frequency. The relative error for the amplitude measurement is less than 5%. As for the 3ω technique, several thicknesses of the GST layer have been considered in order to discriminate the intrinsic thermal conductivity of the GST from the thermal resistance at the interfaces (NiO-GST and GST- Si substrate). The measured amplitude for the fcc-gst is represented in Figure 4 as an example. Figure 5. Schematic representation of the time domain pico second thermoreflectance experiment. Figure 6. Measured and simulated (using the identified thermal effusivity) TDTR for a 400 nm thick fcc-gst layer. Figure 4. Measured amplitude for the fcc-gst sample of different thickness using the PTR. The time domain thermoreflectance (TDTR) is widely used in the field of acoustic and thermal characterization of thin layers at the nano and micro scale [4]. A very schematic description of the experiment implemented at the LMP laboratory (University of Bordeaux, France) is represented in Figure 5. The picoseconds thermoreflectance technique is based on a time resolved pump-probe setup using ultra short laser pulses (795 nm, 100 fs, 5 nj) generated by a Ti:sapphire laser. Radiation of pump is doubled by BBO crystal. The probe pulse is delayed according to the pump pulse up to 12 ns with a temporal precision of a few tens of femtoseconds by means of a variable optical path. The pump beam, for which the optical path length remains constant during the experiment, is modulated at a given frequency of 0.3 MHz by an acousto-optic modulator. To increase the signal to noise ratio, a lock-in amplifier synchronized with the modulation frequency is used. Probe and pump beams have a Gaussian profile and are focused at the surface of the sample at normal incidence by means of objectives. The intensity of the reflected probe beam is measured using a silicon photodiode. The measured reflectivity change at the surface of the sample depends on the temperature rise for both the electron gas and the lattice when time is of the order of the pulse duration (in practice ~0.1 psec). After the thermallization time, the reflectivity change is mainly due to the lattice cooling. In order to increase the thermoreflectance, the GST is capped with an aluminium layer. The GST layer is viewed as a semi infinite medium for the duration of the experiment (1 or 2 ns). A new analytical method has been developed that permit to identify the thermal effusivity of the layer [5]. As represented in Figure 6, an accurate fit between the measured and simulated TDTR is obtained by using the identified value for the effusivity of the GST. 104

106 The 3ω and PTR methods lead to comparable results in terms of the intrinsic thermal conductivity of the GST (~0.2 Wm -1 K -1 for the amorphous phase and ~0.5 Wm -1 K -1 for the fcc-crystalline phase). On the other hand these results are in good agreements with those given by the literature. The effusivity of the GST is provided as additional information from the TDTR technique. Assuming that the specific heat of the layer is that of the bulk, the same value of the thermal conductivity as measured by 3ω and PTR is retrieved. Prospects concern the thermal characterization of GST from the ambient to the melting temperature (~600 C). On the other hand, the same experimental methodology is currently applied to measure the thermal properties of other materials involved in the Non Volatile Memory cell. This work was supported by the European Project TCAMMD (FP6-Marie Curie Intra-European Fellowships) and the MDM- STMicroelectronics 2005 and 2006 projects. [1] J.-L. Battaglia, C. Wiemer and M. Fanciulli, An accurate low-frequency model for the 3ω method, accepted in J. Appl. Phys (2007). [2] J.-L. Battaglia, A. Kusiak, M. Bamford and J. C. Batsale, Int. J. Thermal Sciences 45, 1035 (2006). [3] A. Kusiak and J. L. Battaglia, Eur. Phys. J. Appl. Phys. 35, 17 (2006). [4] C. Rossignol, J. M. Ramppnoux, T. Dehoux, S. Dilhaire, B. Audoin, Rev. Sci. Instrum. 77, (2006). [5] J.-L. Battaglia, A. Kusiak and J.-C. Batsale, Heat Transfer, in Press. 3.5 Binary oxides for resistive switching non volatile memories S. Spiga, C. Wiemer, G. Scarel, M. Perego, G. Tallarida, S. Ferrari, A. Cappella, H. Lu, M. Fanciulli CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy Recently, resistive switching random access nonvolatile memories (ReRAM) are receiving an increasing attention for future technology nodes. Their structure is based on two terminal electrodes that sandwich a resistive change material. Among the materials investigated for these switching memories (organics materials, ternary/quaternary and binary oxides), binary oxides are receiving an increasing interest as: (i) they might be integrated in the current memory devices easier than organic materials and (ii) it is easier to control the stoichiometry of binary than ternary/quaternary compounds. Another important issue for the real implementation of these new materials in memory devices is to find proper electrodes compatible with the standard process flow, such as TiN, W, Cu, Ti. It is therefore necessary to find a proper combination of electrode/resistive switching material. Table I summarizes the binary oxides currently investigated [1-12]. Binary Deposition Electrodes oxides technique Nb 2 O 5 PLD [1] Pt ZrO x PLD [2] Pt ZrO x Reactive Pt sputtering [3] MgO x RF magnetron Pt and Cu sputtering [4] NiO Dc magnetron Pt, Au, Al sputtering [5-7] NiO PLD [8] Pt NiO RF sputtering Pt + oxidation [9] TiO 2 MOCVD, Pt ALD [10,11] TiO 2 Sputtering [12] Pt and Ti Table I. List of binary oxides currently investigated in the literature. In this work we characterize the structural and chemical properties of two promising candidates as switching materials, NiO and TiO 2. Preliminary data on the switching of NiO are also reported. NiO films are grown using two different techniques: atomic layer deposition (ALD) and electron beam deposition. NiO films (10-30 nm thick) are deposited on Si(100) by ALD in the ASM F-120 reactor using the Ni(C 5 H 5 ) 2 and precursor combination at various temperatures (200 C, 250 C, 300 C). Based on the results of structural/chemical analyses performed using XRD, XRR and ToF-SIMS, the growth temperature of 300 C was selected to be the best one. These NiO films are crystallized as grown in the cubic phase (Figure 1), they are dense (electronic density of 1.83 e - /Å 3 ; close to the expected value for NiO) and uniform. ToF-SIMS profiles (Figure 2) evidenced homogeneous NiO profiles through the film thickness. The impurity content (especially C) is low, below 0.1% range. From XPS data, two compositions can be clearly identified: 105

107 on the preferential orientation of the NiO layer. The characterization of films grown on other bottom electrodes is still in progress, aiming at characterizing the oxide/electrode interface and the influence of electrode on film properties. Figure 1. Grazing incidence XRD spectra of 14 nm thick NiO films grown by ALD using the Ni(C 5 H 5 ) 2 + precursor combination. Electron beam deposition of nm thick NiO films is performed starting from NiO pellets (3-6 mm) (vacuum: mbar); the growth temperature is around 40 C. The NiO films are crystallized in the cubic phase and homogeneous (the ToF-SIMS NiO profiles are homogeneous through the film thickness). The measured electronic density is 1.83±0.05 e - /Å 3, comparable to the one measured for ALD grown NiO films at 300 C. The impurity content measured by ToF-SIMS is low and comparable to the one measured 10 0 Intensity (arb.units) Si NiO C Si Time (s) Figure 2. ToF-SIMS depth profiles acquired for 14 nm thick NiO film grown on Si. Figure 3. XPS high resolution O 1s spectra of nm thick NiO films grown by ALD (green) and e-beam (red). NiO which is the predominant one, and Ni(OH) 2 (Figure 3). The latter component is close to the film surface, as evidenced by angle-resolved XPS measurements. A second set of NiO films was therefore grown by ALD (Ni(C 5 H 5 ) 2 and precursors) on metal bottom electrodes (Pt, W and Ni) at the selected temperature of 300 C. Results for films grown on Pt at 300 C (XRD, XRR) evidenced that films are as dense as those grown on Si, and crystallized in the cubic phase. Moreover XRD results evidence an influence of the Pt substrate for ALD grown NiO films (Ni(C 5 H 5 ) 2 + precursor combination). The predominant composition (from XPS) is NiO (Figure 3), Ni(OH) 2 is also detected close to the film surface. A first set of simple devices (flat bottom electrode/ NiO/patterned top electrode) was fabricated to test the switching properties of the NiO films. Preliminary results for 150 nm thick NiO films grown on Pt by electron beam deposition are presented in Figure 4. TiO 2 films were grown by ALD at 295 C using Ti[OCH(CH 3 ) 2 ] 4 as metal precursor and H 2 O or 106

108 Figure 4. Typical I-V characteristics of Pt/150 nm NiO/Pt devices (contact area: 7.8x10-4 cm 2 ). [1] H. Sim, D. Choi, D. Lee, M. Hasan, C. B. Samantaray, H. Hwang, IEEE Electron Device Lett. 26, 292 (2005). [2] S. H. Kim, I. S. Byun, I. R. Hwang, J.-S. Kim, J. S. Choi, S. H. Kim, S. H. Jeon, S. H. Hong, J. H. Lee, B. H. Park, S. Seo, M. J. Lee, D. H. Seo, Y. S. Joung, D.-S. Suh, J. E. Lee and I. K. Yoo, J. of the Korean Phys. Soc. 47, S247 (2005). [3] D. Lee, H. Choi, H. Sim, D. Choi, H. Hwang, M. J. Lee, S. Ae Seo, I. K. Yoo, IEEE Electron Device Lett. (2005). [4] K. W. Jeong, Y. Ho Do, K. S. Yoon, C. O. Kim and J. P. Hong, J. of the Korean Phys. Soc. 48,1501(2006). [5] S. Seo, M. J. Lee, D. H. Seo, E. J. Jeoung, D. S. Suh, Y. S. Joung, and I. K. Yoo, I. R. Hwang, S. H. Kim, I. S. Byun, J. S. Kim, J. S. Choi, B. H. Park, Appl. Phys. Lett. 85, 5655 (2004). [6] S. Seo, M. J. Lee, D. C. Kim, S. E. Ahn, B.-H Park, Y. S. Kim, I. K. Yoo, I. S. Byun, I. R. Hwang, S. H. Kim, J. S. Kim, J. S. Choi, J. H. Lee, S. H. Jeon, S. H. Hong, and B. H. Park, Appl. Phys. Lett. 87, (2005). [7] D. C. Kim, M. J. Lee, S. E. Ahn, S. Seo, J. C. Park, and I. K. Yoo, I. G. Baek, H. J. Kim, E. K. Yim, J. E. Lee, S. O. Park, H. S. Kim, U-In Chung, J. T. Moon, and B. I. Ryu, Appl. Phys. Lett. 88, (2006). [8] S. E. Moon, J. H. Park, E. K. Kim, M. H. Kwak, H. C. Ryu, Y.T. Kim, S. J. Lee, S. Maeng, K. H. Park, K.-Y. Kang, B. H. Park J. of the Korean Phys. Soc. 49, 1066 (2006). [9] K. Kinoshita, T. Tamura, M. Aoki, Y. Sugiyama, H. Tanakaet, Appl. Phys. Lett. 89, (2006). [10] C. Rohde, B. J. Choi, D. S. Jeong, S. Choi, J. S. Zhao, C. S. Hwang, Appl. Phys. Lett. 86, (2005). [11] B. J. Choi, D. S. Jeong, and S. K. Kim, C. Rohde, S. Choi, J. H. Oh, H. J. Kim, C. S. Hwanga, K. Szot, R. Waser, B. Reichenberg, S. Tiedke, J. Appl. Phys. 98, (2005). [12] Y. H. Do, K. W. Jeong, C. Ok Kim, J. P. Hong, J. of the Korean Phys. Soc. 48, 1492 (2006). Figure 5. Grazing incidence XRD spectra of TiO 2 films grown on Si by ALD using (red, thickness 16.5 nm) and H 2 O (green, thickness 13 nm) as oxygen precursors. as oxygen sources. Si, Pt and Au electrodes are used as substrates. As grown films are crystallized in the anastase phase of TiO 2, and no significant variation is observed as a function of oxygen precursor (Figure 5). Moreover, no influence of the substrate (Si and Pt) was observed. The electronic density measured by XRR is 1.01 ± 0.05 e - /Å 3 for both H 2 O and grown TiO 2, lower than the expected one for TiO 2 polymorphs ( e - /Å 3 ). The surface roughness, measured by AFM, is close to 10% of film thickness. XPS analyses show that the film composition is TiO 2. In summary, we grew polycrystalline NiO and TiO 2 films on Si and metal substrates to be used as resisitive switching materials. Preliminary data on the electrical properties are obtained for NiO grown by e-beam. This work was partially supported by the European project EMMA Emerging Materials for Mass-storage Architectures (IST-FP6) and by the industrial research project with STMicroelectronics. 3.6 Electrical and chemical investigations of ZnO layers grown by Atomic Layer Deposition for low thermal budget non volatile memories S. Ferrari 1, E. Speets 1, N. Huby 1, A. Pirovano 2, E. Guziewicz 3, A. Wójcik 3, M. Pra 4, P. Lugli CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy STMicroelectronics, Agrate Brianza, Italy Institute of Physics, Polish Academy of Sciences, Warszawa, Poland Institute for Nanoelectronics Technical University of Munich, München, Germany A number of alternative non-volatile memory concepts exploiting various electrical and magnetic properties of different functional materials are currently being 107

109 Intensity (arb. u.) Intensity (arb. u.) Sputter time (arb. u.) A B Zn(COO)2 + H2O ZnO Si Zn(CH2CH3)2 + H2O ZnO H C OH Si ZnO H C OH Si ZnO Sputter time (arb. u.) Figure 1. TOF- SIMS depth profile of ZnO deposited on silicon. Inside the ZnO layer can be identified 2 domains. Domain A (0-200s) and B (200s-1000s). investigated. Among them the so called Resistive Random Access Memory (RRAM), based on the use of materials able to switch between two states with different resistance, are considered with great interest. The search for appropriate diodes to be integrated as selection elements into RRAM based devices arranged in cross-bar architectures is of great interest. We are investigating the use of diodes based on ZnO as possible selection elements for cross bar memory application. ZnO shows high mobility even in its amorphous state, can be deposited at low temperature and in the back end of the line. We investigated chemical and electrical properties of ZnO layer grown by Atomic Layer Deposition at low temperature (below 200 ), for application as diodes in crossbar memory devices. We have investigated the chemical and electrical properties of ZnO films grown by Atomic Layer Deposition (ALD). We considered two different types of ALD processes. The first one uses Zn(COO) 2 and H 2 O as zinc and oxygen precursors respectively, Si while the second is using Zn(CH 3 CH 2 ) 2 and H 2 O as zinc and oxygen precursor. The use of the di-etyl zinc allows much lower growth temperatures given the high reactivity of such compound. First of all, the time-offlight Secondary Ions Mass Spectrometry (ToF-SIMS) was used to characterize the chemical composition of the ZnO films. The film grown from zinc acetate shows a high content of carbon and hydrogen as shown in Figure 1. Such contaminants are particulary high in the region close to the surface (domain A), while the use of di-etyl zinc allows to obtain higher quality films. For these reasons, we focus on thin films grown by the Savannah reactor since the electrical properties will be connected to the quality of the layer. We measured the electrical properties of the ZnO by fabricating transistors and arrays of junctions. We first realized thin film transistors with ZnO as the conductive channel. The transistor allows to estimate the intrinsic electrical properties of the ZnO, like the mobility, the carrier density that are needed in order to optimize the semiconductor for the junctions. The first measurements were performed with ZnO grown at 170 C. They did not show any field effect, according to the high conductivity of the semiconductor. Indeed, Hall measurements have highlighted a conductivity σ = 259 (Ω.cm) -1 and a carrier density n = cm -3. Simulations of the transistor characteristics based on a drift diffusion model show that a carrier concentration higher than cm -3 will considerably decrease the electrical performances of the device. Using the same precursor and lowering the growth temperature at 100 C enable us to lower the carrier concentration to cm -3. The devices based on this Figure 2. Transistor output curves. W/L=5mm / 12 μm. The inset shows a picture of the interdigited transistor. 108

110 ZnO layer showed good field effect properties as presented on the Figure 2. The saturation mobility, obtained from plotting the square root of the saturation drain versus gate voltage was 1 cm 2 /Vs. The low on/off ratio is the sign of a still high carrier density in the ZnO semiconductor. It is well known that the conductivity in n-type ZnO is due mostly to oxygen vacancies [1], improvements in the on/off ratio are expected upon decreasing the oxygen vacancies, thus increasing the oxygen content in the ZnO films. We are currently investigating the use of post deposition treatements such as annealing [2] or plasma treatments [3]. Figure 5. I-V curves (simulations) as a function of doping concentration. The four curves correspond to 10 16, and cm -3. It has been confirmed by simulation that the carrier concentration should be below cm -3 in order to achieve an on/off ratio of 10 4 required for non-volatile memory application as shown in Figure 5. This work was supported by the European Project VERSATILE. Figure 3. Picture of a cross bar memory device, with 12 μm 12 μm and 3 μm 3 μm area junction on the left and right respectively. Array of Shottky junctions made by Al / ZnO / Au stacks were fabricated by conventional Photolithography. An example picture of the array is shown in Figure 3. The ZnO was patterned in order to obtain dots at the cross points of the metal lines with sizes of 25x25 down to 3x3 microns. Figure 4. J-V characteristics of a 6 6 μm junction based on the structure Al / ZnO / Au. The Figure 4 shows current density voltage curves for a junction Al / ZnO / Au. The junction shows an asymmetric shape demonstrating its rectification properties. The high carrier concentration leads to a relatively low on/off ratio. [1] M. Godlewski, E. M. Goldys, M. R. Phillips, R. Langer and A. Barski, J. Mater. Res 15, 495 (2000). [2] M. S. Aida, E. Tomasella, J. Cellier, M. Jacquet, N. Bouhssira, S. Abed and A. Mosbah, Thin Solid Film 515, 1494 (2006). [3] H. L. Mosbacker, Y. M. Strzhemechny, B. D. White, P. E. Smith, D. C. Look, D. C. Reynolds, C. W. Litton, L. J. Brillson, Appl. Phys. Lett. 87, (2005). 109

111 4. Spintronics 4.1 Atomic Layer Deposition of complex magnetic oxide 4.2 Characterization of Fe/high-κ oxide interfaces 4.3 Shallow donor electron spin coherence and manipulation in Si and SiGe 4.4 Study of the static magnetic field and microwave irradiation response of the random telegraph signal in MOSFETs for qubit implementation 4.5 On-line and off-line Mössbauer spectroscopy investigation of the magnetic properties of oxides Conventional electronics relies on the charges of electrons and holes. Among long-term alternative concepts investigated to reduce device size and power consumption, the possibility of using the spins of electrons to gain new functionalities is very intriguing. Several device concepts can be envisaged based on spin, covering a range of complexity and maturity. In particular, magnetic sensors based on spin-valve giant magnetoresistance (GMR) multilayer, and magnetic random access memory (MRAM) based on magnetic tunnel junction (MTJ) have attracted great interest for spin-based electronics applications. In both GMR and MTJ devices, the spinpolarized current flow is manipulated by a magnetic field controlling the orientation of magnetic moments in the magnetic thin films. A MTJ consists of two ferromagnetic (FM) layers, i.e. Co and Fe, separated by an ultrathin layer of insulator, typically an oxide, with a thickness of about 1 nm. The insulating layer is so thin that electrons can tunnel through the barrier if a bias voltage is applied between the two metal electrodes. In MTJs the tunneling current depends on the relative orientation of the magnetization of the two FM layers, which can be changed by an applied magnetic field. This phenomenon is called tunneling magnetoresistance (TMR). Nowadays MTJs that are based on transition-metal ferromagnets and Al 2 barriers can be fabricated with reproducible characteristics and with TMR values up to 70% at room temperature. Recently large values (> 200 % at room temperature) of TMR observed in epitaxial MTJs based on Fe and MgO further boosted interest in spin dependent tunneling. The interface between the FM electrodes and the tunnel barrier is crucial in determining the TMR. The interfacial density of states directly affects the tunneling spin polarization, and a deep knowledge of the structural and magnetic properties of the interface is required to better engineer the MTJs. Current efforts in designing and manufacturing spintronic devices involve two different approaches. The first is perfecting the GMR and MTJ structures 110

112 by either developing new materials with larger spin polarization of electrons or making improvements or variations in the existing devices that allow better spin filtering. The second effort focuses on finding novel ways of both generation and utilization of spin-polarized currents. These include investigation of spin transport in semiconductors and looking for ways in which semiconductors can function as spin polarizers and spin valves. Semiconductorbased spintronic devices present some technical advantages over those based on metals and may serve, in general, as multifunctional devices. In addition it would be much easier for semiconductor-based devices to be integrated with traditional semiconductor technology. While there are clear advantages for introducing semiconductors in novel spintronic applications, many basic questions related to the combination of semiconductors with other materials to produce a viable spintronic technology remain open. The spin transport across the interface formed between a semiconductor and another material (metal, insulator, semiconductor) is far from well-understood. In the past, one of the strategies to advance understanding of spin transport in hybrid semiconductor structures was to directly borrow knowledge obtained from studies of more traditional magnetic materials. However, there is also an alternative approach involving the direct investigation of spin transport in all-semiconductor device geometries. In such a scenario the recently discovered Ga 1-x Mn x As, Ge 1-x Mn x, ZnO:Mn type ferromagnetic materials where Mn impurities act as dopants could be employed to tailor spin transport properties. Intriguing reports of magnetism in undoped oxides such as HfO 2, possibly related to intrinsic defects, are also opening new scenarios. Nuclear and electronic spins provide perfect candidates for quantum bits (qubits) as their Hilbert spaces are generally well-defined and their decoherence relatively slow. The development of silicon-based qubits not only may take advantage of all the ingenuities developed by modern nanoelectronics, but also aims at a scalable technology. Currently the spintronic research effort at the CNR-INFM MDM National Laboratory focuses on the following directions: Development of all-oxide MTJs based on ALD deposition processes. Development of MTJs based on ALD growth of the ferromagnetic metal (Fe, Co) and of the tunnel oxide. Structural characterization of the FM/Insulator (oxides) interface, where FM is a Fe-based electrode. Investigation of magnetism in semiconducting and insulating oxides (ZnO, HfO 2, Lu 2, ). Qubits based on electron spins of shallow donors and other defects in silicon. Electron spin decay, coherence, detection, and manipulation in semiconductors. Some of the activities presented in this chapter are carried out with the partial support of the Fondazione Cariplo, SOLARIS project. In the following the main research activities carried out at MDM in the spintronics field will be described. In addition to the MTJs and semiconductorbased spin devices where the information is manipulated and stored classically (bit), a long-term and ambitious subfield of spintronics is the application of electron and nuclear spins to quantum information processing and quantum computation. 111

113 4.1 Atomic Layer Deposition of complex magnetic oxides M. Georgieva 1, H. Lu 1, M. Perego 1, G. Scarel 1, A. Zenkevich 2, and M. Fanciulli 1 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy 12 Moscow Engineering Physics Institute, Moscow, Russia Atomic layer deposition (ALD) is a very successful technique for the deposition of high quality and conformal oxide films. At the MDM laboratory a great number of different oxide materials have been grown using this technique. Among these are MnO, LaAlO, LuMnO, LuLaO, MgO and several other high-κ oxides (see Chapter 1 and 2) In the recent years, the magnetic oxide materials of the perovskites type La 1-x +3 A x +2 Mn, where A is a divalent alkaline metal (Sr,Ca,Ba,Pb etc..) have been studied extensively due to the newly discovered and technologically important properties such as colossal magnetoresistance (CMR) and their potential application in the field of Spintronics [1]. Figure 1 illustrates the perovskites unit cell. and metallic properties appear almost simultaneously. It is believed that the double exchange interaction between the Mn 3+ and Mn 4+ pairs control the magnetic and electrical properties of these materials. These magnetic oxides have a high degree of spin polarization and therefore are expected to have enhanced spin-dependent transport properties. In fact, a full spin polarization of 100% is theoretically predicted for the so called half-metallic materials. The electronic density of states is completely spin polarized at the Fermi level and the conductivity is dominated by metallic single-spin charge carriers. This offers potential technological applications as a singlespin electron source and magnetic sensors. It is shown experimentally that the La 0.7 Sr 0.3 Mn belongs to this class of materials [4]. Another field of application of these materials is as ferromagnetic electrodes in Magnetic Tunnel Junction devices (MTJ) [5]. Most of the work in this field focuses on La 1-x Ca x Mn (LCMO) and La 1-x Sr x Mn (LSMO) systems as they have high T C (above RT) and high magnetoresistance values. The phase diagram of La 1-x Sr x Mn compound is shown in Figure 2. It is clear that this material is magnetic in the range between x = 0.2 to 0.45 with the best magnetic properties for x=0.33 (Tc~ 370K). The magnetic and electrical properties of the manganites are very much affected by compositional and microstructural variations caused by stress or strain [6]. Therefore the deposition on to nearly latticematched substrates is most desirable. The best substrates Figure 1. Perovskite unit cell structure of La 1-x +3 A x +2 Mn. The pseudocubic lattice parameter, a, varies between Å depending on the substitution metal A. The crystal structure can vary with regard to the composition [2]. The CMR effect in these systems appears close to the transition temperature (Curie temperature, T C ) from paramagnetic insulating state to ferromagnetic metallic state [3]. The ferromagnetic Figure 2. Phase diagram of La 1-x Sr x Mn [1]. T C is the Curie temperature and T N is the Néel temperature. PM,PI,FM,FI,AFM and CI denote paramagnetic metal, paramagnetic insulator, ferromagnetic metal, ferromagnetic insulator, antiferromagnetic metal and spin-canted insulator states, respectively. 112

114 Figure 3. XRR measurements -the colored lines are the experimental data at the different growth temperatures and the black lines are the corresponding fits- a) and XRD spectra b). used up to now are SrTi, LaAl and NdGa. Another newly developed crystal with the perovskite structure is the (LaAl ) 0.3 (Sr 2 AlTaO 6 ) 0.7 (LSAT). The mismatch between the LSAT (lattice parameter 3.87Å) and the La 0.67 Sr 0.33 Mn (pseudocubic lattice parameter 3.889Å) is only -0.49%. In our experiments we have used Si/SiO 2 substrates, but we are planning to use STO, LAO and LSAT for the growth of LSMO films. The most common techniques for the LSMO films depositions are Molecular Beam Epitaxy (MBE), Pulsed Laser Deposition (PLD) and Sputtering [7,8] from targets with the exact desired composition. All these techniques have been successful in producing good quality films, although they suffer from problems with limited substrate sizes, insufficient compositional control over larger areas, high cost etc. We believe that the ALD could give an alternative to the above mentioned deposition techniques combined with a low cost and the ability to deposit material conformally over large areas and with low compositional variations. To deposit the manganites films by means of ALD three different metal precursors and an oxygen source should be used in order to vary the stoichiometry and obtain the desirable composition. By following the work of Nilsen [9] done on LaCaMnO films, first we have deposited LaSrMnO films on to Si/SiO 2 (10nm) substrates at two different growth temperatures: 283 C and 450 C. The films have been grown in the ALD F-120 ASM-Microchemistry Ltd reactor using the corresponding metal (thd) 3 complex precursors and Growth temperature Composition, RBS Fit parameters, XRR LaSrMnO t, Å 350 La 0.7 Mn 1 Sr (550Å) q c ~ q c, Å Tg=283 C La: cm -2, Mn: σ, Å 9 cm -2, Sr: cm -2, O: cm -2 SiO2 t 83 q c ~ q c σ 4 LaSrMnO t, Å 424 La 0.8 Mn 0.6 Sr 0.4 [Si 0.4] (600Å) q c ~ q c, Å σ, Å 25 LaSrMnO t 417 Tg=450 C La 0.7 Mn 1.2 Sr 0.3 [Si 0.4] (700Å) q c ~ σ 71 SiO2 t, Å 84 q c ~ q c, Å σ, Å 6 Table I. RBS obtained composition and XRR fit parameters film thickness (t), critical edge (q c ) and surface roughness (σ) for the samples grown at 283 C and 450 C. 113

115 ozone as an oxygen source. The film growth sequence was as follows: [2 cy. La(C 11 H 19 O 2 ) 3 (183 C) + (18 C) ( ), 3 cy. Mn(C 11 H 19 O 2 ) 3 (T s = 133 C) + (18 C) ( ), 1 cy. Sr (C 11 H 19 O 2 ) 2 (T s = 225 C) + ( )s, 2 cy. La(C 11 H 19 O 2 ) 3 (183 C) + (18 C) ( ), and 4 cy. Mn(C 11 H 19 O 2 ) 3 (T s = 133 C) + (18 C) ( )]. The films were characterized by RBS, ToF SIMS, XRR and XRD. The results on the films composition given by the RBS measurements are shown in table I. It can be seen that the film grown at low temperature has a composition very close to the desired one (La 0.7 Sr 0.3 Mn ). The film grown at 450 C is not homogeneous, since the data could not be fitted just with one single layer. This result was also confirmed by the ToF SIMS (not shown here) and the XRR measurements. Probably at higher temperatures the different precursor s decomposition rate is changing and the growth becomes more difficult to control. There is also some Si content in this film, which is most likely caused by the intermixing between the film and the substrate at high temperature. The XRR spectra are shown in Figure 3a, and the fits parameters are given in table I. From the fit parameters, one can see that the films grown at different temperatures are different in thickness and roughnessthe film grown at Tg = 283 C is much smoother and thinner, compared to the one grown at Tg = 450 C. The critical edge (q c ) is ~ and it is lower than the expected critical edge ~0.049 for the La 0.7 Sr 0.3 Mn compound. Correspondingly the electronic density of our films is lower than that expected for the La 0.7 Sr 0.3 Mn perovskite type films. This could be partly explained by the fact that both films have a very low degree of crystallization and are mainly amorphous, as seen from the XRD spectra shown in Figure 3b. Never the less the preliminary results we have obtained in regard to the composition of the low temperature ALD grown LSMO film are very encouraging. We expect that by using a good lattice matched single crystal substrates we can improve the as grown films crystallinity in the low temperature range C. The ultimate aim of this study is to utilize the ALD grown LSMO films as FM electrodes in MTJ devices realized by means of photolithography and wet etching processes. Additional work is in progress on the CVDlike deposition of Co and Fe thin magnetic films in order to be used as alternative ferromagnetic electrodes in the MTJ stack. This work was partly supported by the CARIPLO foundation through the SOLARIS Project ( ). [1] E. Dagotto et al., Physics Reports 344 (2001) [2] A. G. Belous, O. I. V yunov, E. V. Pashkova, O. Z. Yanchevskii, A. I. Tovstolytkin and A. M. Pogorelyi, Inorganic Materials, 39 (2003), [3] J. W. Lynn, R. W. Erwin, J. A. Borchers, Q. Huang, A. Santoror, J-LPeng, Z. Y. Li, Phys. Rev. Lett. 76 (1996) [4] J.-H. Park, E. Vescovo, H. J. Kim, C. Kwon, R. Ramesh, and T. Venkatesan, Letters to Nature, 392 (1998), 794. [5] Yu Lu, X. W. Li, G. O. Gong, Gang Xiao, A. Gupta, P. Lecoeur, J. Z. Sun, Y. Y. Wang and V. P. Dravid, Physical Review B, 54 (1996), R8357. [6] C. Kwon, M. C. Robson, K.-C. Kim, J. Y. Gu, S. E. Lofland, S. M. Bhagat, Z. Trajanovic, M. Rajeswari, T. Venkatesan, A. R. Kratz, R. D. Gomez, R. Ramesh, Journal of Magnetism and Magnetic Materials 172 (1997), [7] J. M. Coey, M. Viret, S. von Molnar, Adv. Phys. 48 (1999) 167. [8] A. M. Haghiri-Gosnet, J. P. Renard, J.Phys.D: Appl. Phys. 36 (2003) R127. [9] Ola Nilsen, Martin Lie, Helmer Fjellvag, and Arne Kjekshus Growth of Oxides with complex stoichiometry by ALD technique, exemplified by growth of La 1-x Ca x Mn in Rare earth oxide thin film: growth, characterization, and applications, Edited by. M. Fanciulli and G. Scarel, Springer-Verlag, (2007). 4.2 Characterization of Fe/high-κ oxide interfaces R. Mantovan 1, C. Wiemer 1, A. Zenkevich 2, and M. Fanciulli 1 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy 12 Moscow Engineering Physics Institute, Moscow, Russia The interfaces between ferromagnetic layers and tunnel oxides play a crucial role in the performances of novel spin-based electronic devices such as ferromagnetic tunnel junctions, which have attracted considerable interest due to their potential applications in Magnetic-RAM [1]. The achievement of an ideal flat-interface is of paramount importance in order to enhance the performances of such devices. Here, we report on the structural and magnetic properties 114

116 of the Fe/Al 2, Fe/ZrO 2, Fe/HfO 2, and Fe/Lu 2 interfaces, as obtained by performing conversion electron Mössbauer spectroscopy (CEMS). We estimate with CEMS the reactivity of Fe atoms at the different interfaces (intermixing). Complementary information are obtained by X-ray diffraction (XRD) and X-ray reflectivity (XRR). The Al 2, ZrO 2, HfO 2, and Lu 2 oxides (in the 7 to 20 nm thickness range) are deposited by atomic layer deposition (ALD) on a Si(100) substrate. Their surface roughness, as measured by atomic force microscopy (AFM), is 0.1, 1.1, 0.6, and 2.4 nm respectively. In order to investigate the interface with CEMS, a thin layer (<2 nm) of 57 Fe is deposited on top of the oxides by pulsed laser deposition (PLD) at room temperature (RT), followed by the deposition of nm of 54 Fe. The total Fe thickness is between 15 to 23 nm. CEMS measurements have been performed at RT using a 50 mci 57 Co source in a Rh-matrix, which is moved by a standard constant acceleration drive. The samples are incorporated as electrodes in a parallel-plate avalanche counter. Isomer shifts are given relative to α-fe. Grazing angle XRD and XRR measurements have been performed on Fe/ZrO 2, Fe/HfO 2, and Fe/Lu 2. Figure 1 shows the CEMS results for the Fe/HfO 2 and Fe/ Lu 2 samples, which show the least and most reactive interfaces respectively. The CEM-spectra are fitted with the least-squares fitting program NORMOS 90 [2]. Reai l tv e emissi on 5 % 5 % b Fe/HfO b a c Fe/Lu 2 a c Velocity (mm/s) B (T) hf B (T) hf P(B hf ) P(B hf ) Figure 1. CEMS results at RT for the Fe/HfO 2 and Fe/Lu 2 interfaces. The spectra can be analyzed in terms of a fraction that yields to a magnetically split sextet (component a), a fraction represented by a broad magnetic component (component b), and a paramagnetic doublet c (see text for details). The hyperfine field distributions related to the components b are reported as well. All the CEM-spectra show the presence of both magnetic and paramagnetic phases: a fraction that yields a magnetically split sextet, which is characterized by a magnetic field close to 33 T and denoted with α- Fe (component a in Figure 1); a fraction represented by a distribution of sextets (component b); and a low paramagnetic contribution c (<5%). α-fe is not sensitive to the reactions at the interface, while the other components account for the atomic-mixing at the interfaces [3]. For all the samples, the broad magnetic component b is fitted with a Gaussian distribution of hyperfine fields (B hf ). The corresponding B hf distributions are reported at the right side of the spectra in Figure 1. We proceed with the identification of all the different phases formed at the interfaces, by analysing the CEMS and XRD results. Fe/Al 2 interface. The equilibrium phase diagram of the Fe-Al intermetallic system includes the D - structure-type Fe 3 Al alloy, which is ferromagnetic at RT [4]. In Fe 3 Al, the Fe atoms are distributed between two different sites: one with 8 Fe nearest neighbours (n.n.), and the other with 4 Fe and 4 Al as n.n. We identify the component b at this interface, as corresponding to the Fe 3 Al phase being the distribution due to a possibly varying Al-atoms concentration around the 57 Fe nucleus. The paramagnetic contributions are related to Fe 3+ atoms substituting Al 3+ in the Al 2, and for Fe 2+ in an organometallic biferrocene compound respectively [3]. Fe/ZrO 2 interface. In amorphous alloys, the Fe-Zr system exhibits magnetic ordering below RT for Fe contents close to 90 at.% and the corresponding Mössbauer spectra are characterized by a broad sextet with a peak in the 80 K B hf distribution, located at 22 T [5]. Hydrogenation of these amorphous systems increases both the Curie temperature T C (up to 400 K) and the magnetic moment at the Fe site [5]. The H- absorption from the Fe-Zr (and Fe-Hf) intermetallic compounds is likely to occur due to their largely negative enthalpy of formation, and a relatively stable ternary hydride can be expected. We conclude that the distributions of magnetic sextet at this interface is due to the amorphous Fe x Zr 1-x alloy (x=90), and its hydrogenated phase as well [3]. The H incorporation possibly originates from the two-steps process to grow the samples and in particular to air exposure of the oxides prior to the Fe deposition. The paramagnetic contribution at this interface is thought to be related to a Zr-rich Fe-Zr phase, but its chemical composition has not been determined yet [3]. Fe/HfO 2 interface. The magnetic properties of amorphous Fe x Hf 1-x alloys are very similar to those of 115

117 the corresponding Fe-Zr ones [6]. Compounds with at.% Fe are magnetically ordered at RT having B hf of T (at 4.2 K). Their hydrogenation causes the increase of T C above RT, and B hf rises to 30 T [7]. We identify the B hf distribution as due to the presence of an amorphous Fe-rich Fe x Hf 1-x alloy, together with its hydrogenated phase [3]. This conclusion is supported by the absence, in the XRD spectra, of any signal related to Fe-rich Fe-Hf alloys. On the other hand, XRD detects a small fraction of FeH 2. CEMS results confirm the presence of this phase (component c in Figure 1), and we conclude that a low paramagnetic contribution at this interface is due to FeHf 2. Figure 2. XRD spectra of the Fe/Lu 2 sample. In red are the experimental data. The blue and the black curves respectively indicate the simulations of the experimental data by using the Lu 2 pattern only, and by adding the Fe, LuFe 2, and LuFe 2 H 3 phases. Fe/Lu 2 interface. The crystalline LuFe 2 alloy displays T C =610 K and B hf =18-19 T at RT [8], the latter value being close to that observed for the lower peak in the distribution (Figure 1). From CEMS results, the majority of the Fe atoms are located at sites characterized by higher fields (25-30 T), and this contribution possibly accounts for a hydrogenated LuFe 2 compound, having a higher B hf than LuFe 2 [9]. XRD confirms the presence of crystalline LuFe 2, and reveals a possible contribution from the hydrogenated LuFe 2 H x phase, (Figure 2). XRD does not detect any signal from Lu-rich Fe-Lu compound. The component c in the Fe/Lu 2 CEMspectrum could be related to an amorphous Lu-rich Lu-Fe alloy [3]. We calculate the intermixed fractions at the interfaces (as Fe-equivalent thicknesses) by normalizing the CEMS resonant effects of the paramagnetic+distributi on contributions to the total amount of 57 Fe deposited by PLD. The intermixed fractions for the Fe/HfO 2 and Fe/Al 2 interfaces are 0.5 and 0.7 nm respectively. These values are about half of those occurring at the Fe/ZrO 2 and Fe/Lu 2 ones, being 0.9 and 1.5 respectively [3]. XRR is used to estimate the interfacial layer thickness (including the physical interface roughness) at the Fe/HfO 2, Fe/ZrO 2, and Fe/Lu 2 interfaces, and the values measured are 1, 2, and 3-4 nm, respectively. We observe that the value for the intermixing at the interface, as determined by CEMS, matches the interface layer thickness, as determined by XRR, and the oxide surface roughness as determined by AFM. For oxides showing similar roughness (HfO 2 and Al 2 ), the different observed intermixed volume could be related to different enthalpies of formation for the corresponding Fe-Al and Fe-Hf alloys [3]. To conclude, CEMS was used to investigate with submonolayer resolution the interfaces between Fe and different high-κ oxides. CEM-spectra evidence the presence of both ferromagnetic and paramagnetic phases, the latter giving a low (<5%) contribution to the total area. A considerable contribution to the spectral total intensity is given by broad B hf distributions, which are due to Fe-rich Fe-X binary alloys (X=Al, Zr, Hf, and Lu). Hydrogenation of the Fe-X phases takes place for X=Zr, Hf, and Lu, while at the Fe/Al 2 there is a possible contamination. Contaminations of the oxides surface can originate from their exposure to the atmospheric air prior to the Fe deposition. CEMS and XRD results show that the Fe/HfO 2 and Fe/Al 2 interfaces are the least reactive. A higher intermixing, observed by CEMS, correlates with a higher interfacial layer thickness, determined by XRR, and a larger oxide surface roughness, measured by AFM. We acknowledge Dr. G. Tallarida for the AFM measurements and Dr. G. Scarel for the ALD growth of the high-κ oxides. [1] G. A. Prinz. Science 282, 1660 (1998). [2] R. A. Brand, Wissenschaftliche Elektronik GmBH, Starnberg, Germany (1994). [3] R. Mantovan, C. Wiemer, A. Zenkevich, and M. Fanciulli, Hyperfine Interactions, 169, 1349 (2006). [4] R. Checchetto, C. Tosello, A. Mioltello, and G. Principi, J. Phys.: Condens. Matter 13, 811 (2001). [5] D. H. Ryan and J. M. Coey Phys. Rev. B 35, 8630 (1987). [6] M. C. Lin, R. G. Barnes, and D. R. Torgeson, Phys. Rev. B 24, 3712 (1981). [7] D. H. Ryan, J. M. D. Coey, and J. O. Ström-Olsen, J. Magn. Magn. Mat. 67,148 (1987). [8] U. Atzmony and M. P. Dariel, Phys. Rev. B 10, 2060 (1974). [9] K. H. J. Buschow, P. C. P. Bouten, and A. R. Miadema, Rep. Prog. Phys. 45,937 (1982). 116

118 4.3 Shallow donor electron spin coherence and manipulation in Si and SiGe M. Fanciulli 1, A. Ponti 2, and A. Ferretti CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy CNR Istituto di Scienze e Tecnologie Molecolari, Milano, Italy In recent years, the application of quantum mechanics to solid-state physics has produced a large number of concepts for semiconductor devices with novel functionalities. Quantum mechanical properties start to dominate as device dimensions get smaller. While this is a problem for traditional electronic devices, it is also a prerequisite for the development of quantum technology in general and quantum information processing (QIP) in particular. These possibilities have inspired a number of proposals for creating qubits, gates and quantum registers from semiconductor devices and using them for implementing quantum algorithms. Theoretical studies of the device properties have been undertaken, but the challenging technological hurdles have slowed down experimental progress. Silicon is the main material in mainstream nanoelectronic. However, the mature and advanced silicon technology is facing the challenges of the continuing down-scaling. Nodes where it might be necessary to introduce alternative materials, such as SiGe, Ge, or GaAs, are on the semiconductor road map horizon. Germanium was used in the early stages of microelectronics, but soon substituted by Si, mainly due to the stable and excellent oxide of the latter. Ge has higher electron and hole mobility compared to silicon. These parameters are important for conventional microelectronic applications, but also for some envisaged quantum information processing devices. Combining Si and Ge introduces additional degrees of freedom that permit fine-tuning of material properties as well as the definition of qubits (by quantum dots) and addressing (by g-factor tuning). Since Si and Ge are both group IV elements, issues of compatibility and processing are limited. Ge is used in SiGe RF devices and to build up the strained silicon CMOS template layers. Some of the properties drawbacks or advantages depending on the specific device concept of Ge are the high Z number, which causes a larger spin-orbit interaction, and the larger natural abundance and the larger nuclear spin of 73 Ge (7.8 %, I=9/2) than 29 Si (4.67 %, I=1/2). Both of these properties tend to shorten the coherence time. Nevertheless, the coherence times in group IV materials have been shown to be orders of magnitude longer than in III-V materials. The activity carried out in this context at the MDM Laboratory during the years focused on the experimental determination of shallow donor electron spin coherence in Si and SiGe using pulse electron spin resonance techniques and on the experimental and theoretical (see section 8) investigation of the shallow donor wave function and its manipulation using external electric fields. Other related objectives, such as read-out, are discussed in the next paragraph. Spins of single donor atoms are attractive candidates for large scale quantum information processing in silicon and SiGe [1, 2]. In several solid state schemes, quantum computation is realized through the manipulation of electron and/or nuclear spins of single donors ( 31 P) in Si or in SiGe heterostructures, double donors ( 126 Te) in Si, or impurities (Na) in Si-MOSFETs. When compared with nuclear spins, electron spins allow higher clock rates and require only one swap (spin to charge) for detection. Coherence times in silicon at low temperatures (< 5 K) are sufficiently long to allow fault-tollerant QIP. For isotopically purified crystals coherence times (T 2 ) of the order of ms have been reported while the time necessary for a π pulse is, at a transition frequency of 9.5 GHz, of the order of 20 ns. Several dephasing mechanisms play an important role, depending on the qubit structure, and part of the research effort has been dedicated to the identification of the optimal qubit design. The identification of defects (impurities) with optimal properties for QC is, for example, a crucial issue which may pave the way for qubit high-t operation. Single qubit manipulation has been pursued via hyperfine interaction or g-matrix tuning, while spin-dependent transport in MOSFETs has been investigated for spin detection (see next paragraph). Spin coherence in Si and SiGe has been investigated by pulse X-band electron spin resonance spectroscopy on bulk samples [3, 4]. Two-pulse electron spin echo experiments on P shallow donors in natural and isotopically pure ( 28 Si) silicon and in SiGe alloy, with different P concentrations, as well as other more complex pulse schemes, such as for example CPMG, were used to address spin-spin relaxation times and mechanisms. Electron spin echo envelope modulation (ESEEM) (Figure 1-2) provided information on the donor electron wave function. The latter has been extensively investigated in Si by the electron-nuclear double resonance (ENDOR) technique. However, due to the low nuclear resonance frequency of Ge, ENDOR has technical difficulties in the investigation of Ge related hyperfine interactions. 117

119 Two-pulse ESE experiments performed as a function of the pulse turning angle allowed us to measure the exponential and spectral-diffusion relaxation times depurated by the additional effects of instantaneous diffusion [3, 4]. Spectral diffusion is an important dephasing mechanism which is significantly reduced in the isotopically purified 28 Si sample. The influence of the second pulse, in any sequence aiming at qubit manipulation, must be taken into consideration as, depending also on the total P concentration, instantaneous diffusion could reduce dramatically Figure 2. Electron spin echo decay and magnitude- FT two-pulse ESEEM spectra of silicon germanium alloy samples. ESE decay: (a) FT spectra (b), expanded view in (c) [5]. Figure 1. Magnitude-FT two-pulse ESEEM spectra of Silicon sample. Contribution to the 29 Si ESEEM from shells A, B, and E is well resolved [3,4]. the phase coherence time. According to these results, isotopically purified samples are necessary to reduce the spectral diffusion contribution and, from a comparison with the literature data, the P shallow donor concentration plays a fundamental role to determine the intrinsic phase memory time of this material. Accurate analyses (Figure 1) of the electron spin echo modulation (ESEEM), observed in the natural silicon sample, provide information on the shallow donor wave function. ESEEM peaks have been attributed to the hyperfine-coupled 29 Si nuclei in the various crystallographic positions on the basis of a spectral fit procedure including instrumental distortions. Similar data have been obtained in SiGe (Figure 2) revealing details of the shallow donor wave function so far not determined with ENDOR techniques. Several attempts to detect manipulation of the electronic wave function via external electric fields by observing the hyperfine interaction have been so far not conclusive due to the presence of leakage current leading to local heating [5]. A new dedicated set-up which should allow the application of high electric fields preventing leakage has been designed and is under construction. [1] B. Kane, Nature 393, 133 (1998). [2] R. Vrijen, E. E. Yablonovitch, K. Wang, H. W. Jiang, A. Balandin, V. Roychowdhury, T. Mor, and D. DiVincenzo, Phys. Rev. A 62, (2000). [3] M. Fanciulli, P. Höfer, and A. Ponti, Physica B , 895 (2003). [4] A. Ferretti, M. Fanciulli, A. Ponti, and A. Schweiger, Phys. Rev. B 72, (2005). [5] M. Fanciulli, 2004 IEEE NTC Quantum Device Technology Workshop, Clarkson University, N.Y., USA, May 17-21, (2004). 118

120 4.4 Study of the static magnetic field and microwave irradiation response of the random telegraph signal in MOSFETs for qubit implementation E. Prati 1, M. Fanciulli 1, G. Ferrari 2, and M. Sampietro 2 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy 12 Dipartimento di elettronica e informazione, Politecnico di Milano, Milano, Italy Random telegraph signal (RTS), consisting of the random switching of the current between two levels in a metal-oxide-semiconductor field effect transistor, has been studied for decades. The quantum theory of the RTS in silicon metal-oxide-semiconductor field-effect transistors (MOSFETs) is based on inelastic capture and emission of charge by a trapping center near the Si/SiO 2 interface. Each tunnel transition of the electron between the trap and the channel is assisted by multiphonon emission and absorption, and reflects the effect of electrostatic Coulomb barriers of the charges present in the device [1]. RTS (Figure 1) has been proposed as a viable way towards single spin detection, a key issue for solid state quantum information processing (QIP) [2-4]. In particular, for QIP schemes in which the electron spins are the qubits, read-out of the spin state can be achieved by monitoring the variation of the RTS in electron spin resonance conditions. In this scheme the MOSFET is operated in a static magnetic field and under microwave irradiation. Hence, the response of a trap at the Si/SiO 2 interface of a MOSFET placed in a static magnetic field and/or irradiated by a microwave field, and the occupation of the trap spin states are of considerable interest. We have widely characterized capture and emission of single electrons in point defects under such operating conditions at cryogenic temperatures (below 1 K). Several related results have also been discovered and characterized, such as the stationary current excited into the MOSFET under microwave irradiation, and giant RTS in standard MOSFETs. Recently we have also reported the cyclostationary initialization of the trap charge state [5]. The effect of the static magnetic field on the RTS has been described in [6], and a model has been proposed to take into account the characteristic time ratio r = τ high / τ low change as a function of the magnetic field, where τ high and τ low are the average characteristic Figure 1. The typical trace of the random telegraph signal in our sample during the measurements. This acquisition has been done at V G = mv and V D = 1.6 mv. Below the trace, the reconstructed signal obtained by off-line data processing. times of the two current levels. However, at very low temperatures and high magnetic fields, significant discrepancies between the experimental data and the proposed model have been observed. We have investigated the RTS parameters as a function of static magnetic fields up to 12 T, both in-plane and perpendicular to the two dimensional electron gas (2DEG) localized in the inversion layer of a Si n- MOSFET [7]. Our experiments are performed at liquid 3 He temperature where the effects due to the static magnetic field are expected and proved to be strong. We proposed a model which, taking into account the triplet state of the trap, explains the high field effects. Figure 2 shows the characteristic time ratio r for a trap observed with V G = mv and V D = 15.7 mv (Trap 1) as a function of an external static magnetic field B paral in the plane of the channel and oriented in the direction perpendicular to the electron current flow as shown in the inset of the Figure 2. The drain current was I D = 37.3 na. The low current state of the RTS corresponds to the physical state of an electron captured by the trap, because the ratio r decreases by increasing V G (therefore lowering the trap energy E T ). In order to determine whether the trap was singly occupied, therefore paramagnetic, and capable of capturing a second electron (type 1 2) or an empty trap capable of capturing an electron (0 1), we have analysed the dependence of r on the magnetic field B. The conclusion was that the trap is a 1 2 type. The same sample was examined also with the external static magnetic field B perp perpendicular to the channel. Figure 3 shows the ratio r for another trap (Trap 2) observed with V G = mv and V D = 1.6 mv. At this gate voltage, the channel electron density is n s = 119

121 8x10 11 cm -2. The current was I D = 35 na. This trap has also been investigated with B as discussed later. In both cases the exponential trend is suppressed at high B field because the triplet state becomes probable as the singlet state and the energy levels cross. The need to irradiate the sample to produce spin flip of the electron trapped into the point defect like in [8] motivates the complete characterization of the microwave field effects on the MOSFET and the RTS. In order to discriminate the relevant effects of the microwave on the device under test, we have studied the evidence of a microwave-induced stationary current in MOSFET devices operated in X-band (9.6 GHz) and Q-band (33.9 GHz) resonant cavities, where the structure of the microwave field has well known components. We detect a stationary current generated by the microwave radiation. This microwave-induced the DC value of the drain current changes linearly with the power [9, 10]. This variation depends on the power of the microwave (P μw ) and on the bias condition of the MOSFET (V GS ). Figure 4 shows the experimental evidence of these variations as a function of P μw for the device irradiated in the X-band cavity. Such an effect has been patented a CMOS compatible microwave. The microwave irradiation has another major effect related to the harmonic potential at the drain, gate and source terminals. We demonstrated that also the emission and capture times of the trap may change as a function of the intensity of the microwave field. We prove that the mechanism leading to the observed changes of τ c and τ e under microwave irradiation is related to the voltage oscillation induced by the microwave [11]. Such a conclusion is based on the direct comparison between the measured changes on Figure 2. The ratio r as a function of B. Because of the power dissipated by the current, the temperature of the electrons involved in the RTS of the trap here considered has been evaluated to be 1.5 K, instead of the nominal cold-finger temperature of 245 mk, by fitting the experimental results (black points) with the equation developed in [5] (line). Inset: the magnetic field is in the plane of the 2DEG and perpendicular to the current channel. stationary current has been examined as a function of microwave power, gate voltage, and drain-source voltage. The experimental results can be successfully reproduced by a model exploiting a non linearity of the MOSFET as a component of a circuit coupled with the electromagnetic field. The presence of a current sensible to the applied microwave power reveals a possible intrinsic source of systematic error on average current measurements like one used in [8], if resonant absorption effects of elements of the device or the environment alter the effective microwave power detected by the MOSFET. This effect is also responsible of power dissipation through the two dimensional gas at liquid He3 temperatures, where single spin resonance is believed to be detectable. When the MOSFET is irradiated with the microwave, Figure 3. The ratio r as a function of B perp. The experimental measurement (black points) has been fitted including the triplet state, obtaining a local temperature 2.3 K, instead of the nominal cold-finger temperature of 245 mk. The RTS disappears between 7.75 T and 9.4 T because of magnetically induced localization. The inset shows the direction of the magnetic field, orthogonal to the sample. the characteristic times as a function of the microwave power and the values calculated assuming that these changes are due to the V DS modulation induced by the microwave. To calculate the expected characteristic time values, we average the transition probabilities by weighting along the V DS values harmonically varying by an amplitude obtained from the microwave induced stationary current measurements. The latter provides the conversion factor between the microwave power and the effective voltage amplitude. The effect, shown in Figure 5, can be fully ascribed to the inevitably present coupling between the microwave field and the conductive loop formed by the MOSFET and the connections towards the sensing amplifier: the microwave field induces a harmonic current on the loop, modulating the source and drain 120

122 voltages of the MOSFET. In circuital representation, this corresponds to adding two AC voltage generators at the drain and at the source of the MOSFET with the same frequency of the microwave field. In static condition the characteristic times τ c and τ e change as a function of V D. Exploiting the change of the field distribution in the proximity of a dipole antenna as a function of the frequency of the microwave, we are able to set a frequency where the coupling of the MOSFET with the microwave field occurs only at the drain. To set such condition we used a microwave frequency of GHz. At room temperature, without microwave field applied and at V G =800 mv, the RTS has a mean capture time τ c ranging monotonically from 3 ms to 20 ms for a drain voltage variation from 200 to 800 mv, while the mean emission time τ e remains almost constant at about 0.7 ms. To summarize, the inevitably present electric loop due to the on-chip and off-chip connections of a MOSFET to the external measuring system is responsible for the RTS variation upon microwave irradiation through the modulation of the MOSFET biasing conditions. The experimental observation of the change of the RTS under microwave irradiation has been reported and the proposed model correctly predicts such a change. Such an effect has to be carefully isolated and minimized whenever attempting to electrically detect single spin resonance via RTS. Otherwise, a spurious microwave absorption in the environment may vary the effective power of the microwave field coupled with the MOSFET and produce a change of the RTS characteristics not related to the trap - responsible for the RTS - driven in spin resonance conditions. Any measured τ c, τ e and dc current change in agreement with the proposed model prediction at a given power absorption has to be regarded as a spurious effect and cannot be ascribed to a single spin resonance phenomenon. A new development of the technique is under investigation, known as pulsed testing. Pulsed electrical measurements should reduce the total energy dissipated in the device, and thus minimizes thermal heating, while pulsed microwave irradiation should enable to control the spin state. Figure 4. Drain current variation in the n-mosfet biased in ohmic regime (V DS = 0 V) as a function of the microwave power (9.6 GHz) at different gate bias voltages. Continuous lines guide the eyes through the experimental points (dots). Figure 5. Experimental mean capture (up triangles) and emission (down triangles) times versus microwave field power compared to the theoretical predictions (continuous lines). Experimental conditions: V G = 800 mv, V D = 500 mv, τ c0 = 14 ms, τ e0 = 0.7 ms, ν =15.26 GHz. The data are normalized to the capture time in absence of microwave irradiation. [1] A. Palma, A. Godoy, J. A. Jimenez-Tejada, J. E. Carceller, J. A. Lopez-Villanueva, Phys. Rev. B, 56, 15, 9565 (1997). [2] R. Vrijen, E. E. Yablonovitch, K. Wang, H. W. Jiang, A. Balandin, V. Roychowdhury, T. Mor, and D. DiVincenzo, Phys. Rev. A 62, (2000). [3] Elzerman, Nature, [4] M. Fanciulli, E. Prati, G. Ferrari, and M. Sampietro, Random Telegraph Signal in Si n-mosfets: a Way Towards Single Spin Resonance Detection, AIP Conference Proceedings Vol. 800, Pag. 125, (2005). [5] E. Prati, M. Fanciulli, in preparation. [6] M. Xiao, I. Martin, and H. W. Jiang, Phys. Rev. Lett. 91, 7, (2003). [7] E. Prati, M. Fanciulli, G. Ferrari, M. Sampietro, Effect of the Triplet State on the Random Telegraph Signal in Si n-mosfets, Phys. Rev. B 74, (2006). [8] M. Xiao, I. Martin, E. Yablonovitch, and H. W. Jiang, Nature, 430, 435 (2004). [9] G. Ferrari, L. Fumagalli, M. Sampietro, E. Prati and M. Fanciulli, CMOS fully compatible microwave detector based on MOSFET operating in resistive regime, IEEE Microwave and Wireless Components Letters, 15, 7, 445 (2005). [10] G. Ferrari, L. Fumagalli, M. Sampietro, E. Prati and M. Fanciulli, dc current modulation in field effect transistors operating under microwave irradiation for quantum read out, J. Appl. Physics, 98, (2005). [11] E. Prati, M. Fanciulli, A. Calderoni, G. Ferrari, M. Sampietro, Microwave Induced Effects on Random Telegraph Signal in Si n-mosfets, cond-mat/ , Physics Letters A, in press. 121

123 4.5 On-line and off-line Mössbauer spectroscopy investigation of the magnetic properties of oxides R. Mantovan 1, M. Fanciulli 1, G. Weyer 2, H.P. Gunnlaugsson 2, D. Naidoo 3, R. Sielemann 4, K. Bharuth- Ram 5, T. Agne CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy Institute of Physics and Astronomy, Aarhus University, Aarhus C, Denmark School of Physics, University of the Witwatersrand, South Africa School of Pure and Applied physics, University of KwaZulu-Natal, Durban, South Africa Hahn-Meitner Institute, Berlin, Germany EP Division, CERN, Geneva 23, Switzerland Finding semiconductors and oxides showing ferromagnetism (FM) at room temperature is a challenging task for solid-state physics, and of fundamental importance for future applications in spin-based electronics (spintronics) [1]. The prediction for an ordering temperature Tc>RT in Mn-doped ZnO [2], started a new field of research with the aim to understand and control the magnetic properties of oxides. The generally applied macroscopic magnetization measurements on thin layers are vulnerable to misinterpretations in case of the presence of small magnetic precipitates in the film or in the substrate material. In general, FM is achieved mostly by doping with dilute 3d transition metal impurities, notably Mn, Fe, and Co (in % concentrations), during growth or by subsequent ion implantation. Recently, the discover of unexpected magnetic properties in undoped HfO 2, which is an insulating oxide having the Hf 4+ cation with empty d-shell, opened new fundamental questions about the origin of FM in oxides [3]. Defects are supposed to be directly involved in the occurring of magnetic ordering in HfO 2, and firstprinciples investigations show that Hf vacancy is the only defect creating a stable magnetic moment in this oxide [4]. On the other hand, another group reported that handling (apparently) undoped HfO 2 samples with stainless-steel tweezers, lead to the same magnetic properties as reported in [3]; while using plastic tweezers, the magnetism was not observed [5]. For ZnO, bound polaron-mediated ferromagnetic coupling of 3d magnetic moments within polaron orbits has been proposed [6]. Zn vacancies as well as Zn interstitials, acting as hole and electron dopants respectively, have been proposed theoretically to promote FM in Mnand Co-doped material [7]. Complexes in Fe-doped ZnO, such as Fe Zn -V Zn, have been also shown to be magnetic [8]. It is evident that the conditions and the origin of the magnetism in such materials are not yet understood. Only few analytical techniques, which can give atomic scale information on the electronic configurations and magnetic properties have been applied, among them Mössbauer spectroscopy using the 57 Fe Mössbauer isotope. We perform on-line Mössbauer spectroscopy at the large-scale facility of ISOLDE-CERN. A beam of radioactive 57 Mn + (T 1/2 = 1.5 min) with 60 kev energy is obtained at ISOLDE, by using proton induced (1.4 GeV) fission in a UC 2 target [9]. The short-lived 57 Mn atoms, β - -decay to the 14.4 kev Mössbauer state of 57 Fe. The 57 Mn + ions are implanted with a fluence of cm -2 in single crystals ZnO (provided by Crystec), in HfO 2 and Lu 2 oxides produced by atomic layer deposition (ALD), and in quartz (Suprasil). All the samples were implanted from room temperature up to 1000 K. The samples are annealed during the lifetime of the 57 Mn state. Lattice defects (10 3 /ion) are produced simultaneously, and their interaction with the implanted probe atoms can be studied by varying the implantation temperature. The local concentration of the Mn/Fe atoms is about cm -3, i.e. a fraction of percent of the implanted species. This excludes the presence of secondary magnetic phases in the oxides. The Mössbauer spectra of the daughter 57 Fe atoms are recorded by a resonance detector, which is mounted Relative emission (arb. units) 462 K 700 K Velocity (mm/s) Figure 1. On-line Mössbauer spectra at 462K and 700 K of a single crystal ZnO sample implanted with radioactive 57 Mn + at ISOLDE-CERN. 122

124 on a conventional Mössbauer drive systems outside the implantation chamber. All the spectra have been analyzed by using the Vinda fitting program [10]. Figure 1 shows the on-line Mössbauer spectra of an Mn-implanted ZnO sample at 462 K and 700 K. The spectrum at 462 K shows 90% spectral area of magnetic components. A well-resolved magnetic sextet (red) can be identified by its four outer lines, implying a large magnetic hyperfine field B hf = 48 T. In addition there is the presence of two broad, asymmetric distributions of sextets with varying, lower B hf values, being 37 T (blu) and 12 T (cyan) respectively. In the central part, a quadrupole-split component (green) as well as an apparent single line (orange) are visible directly. The latter has basically disappeared from the 462 K spectrum, which shows the highest magnetic fraction. At T> 600 K the dominant component is a single line, unambiguously assigned to Fe 2+, which results from the annealing of the sextet as well as from parts of the 37 T distribution. The two distributions are still visible in the spectrum at 700 K. The nature of all the spectral components will be discussed in a forthcoming paper [11]. The majority of the Fe atoms probe a magnetically-ordered local surrounding up to 600 K. From the temperature dependence of the magnetic hyperfine splitting for the sextet an even higher Tc value is indicated. This shows that a dilute magnetic semiconductor with Tc >> RT can be produced by ion implantation. Since the magnetic signals in the spectra were visible already after implantation of only probe atoms, a carrier-mediated ferro- or antiferromagnetic coupling appears unlikely as well as a bound polaron coupling. The evolution of the sextet and the 37 T distribution as a function of the temperature, on the time scale of the 57 Mn lifetime, points rather to the formation (and dissolution at high temperatures) of very small, isolated magnetic clusters involving the substitutional Mn/Fe atoms on Zn sites and at least one vacancy in the neighborhood [11]. The experimental data are in agreement with theoretical predictions [8]. The on-line Mössbauer spectra of the Mn-implanted HfO 2, Lu 2 and quartz are characterized by the presence of both paramagnetic components together with a distributions of magnetic sextets [12]. In particular, the magnetic contributions are observed up to 1000 K, accounting for the % of the total spectral area. In the Mn-implanted Lu 2, the non magnetic contribution progressively decreases with increasing the temperature. This corresponds to a significant increase of the magnetic fraction, i.e. the annealing process seems to favour the magnetic ordering [11]. In contrast to what we observe in ZnO, the average hyperfine field of the magnetic distribution in the Mössbauer spectra of the Mn-implanted HfO 2, Lu 2 and quartz show a constant decrease with increasing the temperature, suggesting a possible different origin for the observed magnetic ordering. In order to better explore and to understand the nature of the magnetism in the Mn/Fe-implanted oxides and semiconductors, off-line experiments are also planned. In particular, conversion electron Mössbauer spectroscopy (CEMS) will be performed in the K temperature range on ZnO single crystals implanted with stable 57 Fe atoms at different energies and doses. When compared to the on-line Mössbauer technique, the off-line CEMS requires a larger concentration of the implanted Fe atoms in order to obtain a measurable effect. The mechanism responsible for the magnetism in the Mn-implanted ZnO, HfO 2, Lu 2, and quartz is currently under study. The large concentration of defects produced during the implantation process, i.e. cation and/or oxygen vacancies or interstitials, could play an important role in the occurring of the magnetic ordering [8]. In the on-line Mössbauer experiments, we produce truly dilute oxide/semiconductors due to the low concentration of implanted atoms. This reasonably excludes that ferromagnetic precipitates are responsible for the observed magnetism. Both online and off-line Mössbauer spectroscopy experiments are planned. The control of the magnetic properties of oxides will give the opportunity to develop a large variety of novel materials for spintronics. [1] S. J. Pearton, W. H. Heo, M. Ivill, D. P. Norton, and T.Steiner, Semicond.Sci. Technol. 19, R59 (2004). [2] T. Dietl, H. Ohno, F. Matsukura, J. Cibert, D. Ferrand, Science 287, 1019 (2000). [3] M. Vankatesan, C. B. Fitzgerald, and J. M. D. Coey, Nature 430, 630 (2004). [4] C.D. Pemmaraju and s. Sanvito, Phys. Rev. Lett. 94, (2005). [5] D.W. Abraham, M. M.Frank,and S. Guha, Appl. Phys. Lett. 87, (2005). [6] J.M.D. Coey et al., Nature Mat. 5, 173 (2005). [7] M.H.F. Sluiter et al., Phys. Rev. Lett. 94, (2005). [8] A. Debernardi and M. Fanciulli, Appl. Phys. Lett., to be published [see chapter 8]. [9] G. Weyer and the ISOLDE Collaboration, Hyperfine Interactions 129, 371 (2000). [10] [11] G. Weyer, H.P. Gunnlaugsson, R. Mantovan, D. Naidoo, R. Sielemann, K. Bharuth-Ram, M. Fanciulli, T. Agne, unpublished. [12] R. Mantovan, Ph.D thesis, Università degli Studi di Milano, Dipartimento di Fisica, January

125 5. Devices based on organic and polymeric semiconductors 5.1 Al 2 gate dielectric in poly (3-hexylthiophene) based transistors 5.2 Atomic Layer Deposited Al 2 as a capping layer for polymer based transistors 5.3 A novel method for the production of top contact thin film transistors based on organic semiconductors Organic electronics (OE) is proposed as a platform for a wide variety of electronic products exploiting the electrical and optical properties of conjugated polymers and molecules. OE showed incredible progresses during the last thirty years through the understanding and tailoring of the electron transport mechanisms in organic materials. Indeed since their discovery in the 80 s the science of OE developed into a branched and highly relevant discipline, with application in electronics, optoelectronics, sensors and more. From the application point of view small Organic Light Emitting Diodes (OLED) are already an established and commercially available technology, while for other applications like organic solar cells big companies are making significant investments. Transistors are the essential component of every electronic device therefore Organic Thin Film Transistors (OTFT) are the subject of intense investigation among different groups worldwide. OTFT are not considered for the replacement of the silicon based microelectronic, but rather for the replacement of the Thin Film Transistor technology that is currently based on amorphous silicon and finds application in fields like back plane for active matrix displays, low performances low cost devices like RFID and more. Among the critical points addressed by the study of OTFT are the understanding of the 124

126 electron transport mechanisms, to improve the OTFT performances and the development of processes that allow the manufacturing and integration of organic semiconductors into real devices. At the MDM Laboratory we are exploiting our competences on the growth of materials by Atomic Layer Deposition (ALD), to fabricate and to study transistors that are coupling Al 2 with polymeric semiconductors, as a way to improve transistor performances and to elaborate new integration strategies for the organic semiconductors. In paragraph 5.1 we study the use of ALD deposited Al 2 as gate dielectric for OTFT. We show that the use of Al 2, with a dielectric constant much higher than the dielectrics commonly used in OTFT, allows a significant decrease of the working voltage of the transistor, although the nature of the oxide/polymer interface is quite peculiar and affects the properties of the OTFT. We show that by proper interface engineering it is possible to achieve remarkable transistor properties. damage the underlying semiconductor, and on the other provides an excellent barrier against device degradation upon atmospheric exposure. Finally paragraph 5.3 shows a strategy for fabricating transistors on top of polymeric semiconductor layers by conventional optical lithography after capping with Al 2. The activities presented in this chapter are the results of two projects both founded by Fondazione Cariplo: TE- SEO Project that run between 2004 and 2005, and PROTEO project started in 2006 and still running in collaboration with CNR ISMAC and the Politecnico di Milano Dipartimento di Elettronica. Paragraph 5.2 shows how the ALD deposited Al 2 can be used as a capping layer for OTFT. Since one of the problems of the organic semiconductors is their sensitivity to the atmospheric exposure. We show a method that allows the deposition of the oxide on top of the organic semiconductors, that on one side does not 125

127 5.1 Al 2 gate dielectric in poly (3-hexylthiophene) based transistors E. Peron 1, F. Perissinotti 1, G. Tallarida 1 and S. Ferrari 1 L. Fumagalli 2, D. Natali 2 and M. Sampietro 2 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy 12 Politecnico di Milano, Dipartimento di Elettronica e Informazione, Milano, Italy Figure 1. XRD spectra of the P3HT thin films on bare and functionalized Al 2 surfaces. The solid lines represent fits to the reflectivity/diffraction data based on refined structural models. The role of the dielectric and of the semiconductor/ dielectric interface in the determination of the performances of the Organic Thin Film Transistor (OTFT) has been realized only recently. Considering that the channel in organic TFT is limited to the first one or two monolayers, it is not surprising that the dipole moment, the templating effect and the charged mobile impurities of the dielectric affects strongly the charge transport process in the semiconductor channel. We considered the use of Al 2 as gate dielectric in bottom contact transistor structures. Since transistors using Al 2 gate dielectrics show in general lower charge carrier mobilities, as compared to those based on the more commonly used SiO 2 we considered the use of different self assembled monolayers (SAM) to improve the transistor mobility in OTFT based on RR-P3HT. We selected two groups of SAMs, alkyl-silanes and alkyl phosphonic acids, which differ for their binding group, and in each group we tested molecules of different chain lengths, going from a single methyl group up to a octadecyl chain. Figure 1 shows X-ray Bragg-Brentano profiles of the P3HT films deposited on the differently functionalized Al 2 surfaces. The films deposited on bare and n- octyl phosphonic acid (NOPA) (and on all short chain SAMs not shown here) coated surfaces do not show any diffraction peak, suggesting that the films are amorphous. The profiles of the P3HT films deposited on octadecyl trichloro silane (OTS) and octadecyl phoshonic acid (OPA) on the other hand show a broad diffraction peak that can be related to the <100> reflection of the P3HT. From the diffraction peaks we can estimate that the interplanar distance in the P3HT crystals are 18.0 Å for OPA and 17.6 Å for OTS. The appearance of the <100> diffraction peak in the outof-plane direction is therefore specific of films spun cast on long chain SAM. We explain this phenomenon considering that the long chain SAMs can achive an high degree of order that provides a templating effect on the spun cast P3HT. Indeed the hexyl chains of the P3HT can allign themself in a ordered way by interacting with the ordered SAM monolayer. The transistors using the different SAMs on Al 2 and SiO 2 have been electrically characterized by measuring their transfer and output characteristic in the saturation regime. Figure 2. Mobility values calculated at V G =-15 V on transistors using Al 2 dielectrics (light gray) and SiO 2 (dark gray) dielectric, given for reference. Figure 2 summarizes the average mobility achieved on bare and functionalized Al 2 measured at V G =-15V. With respect to the bare Al 2, which has a mobility around a few 10-4 cm 2 /Vs, the use of short chain SAM does not improve the mobility. A sizeable boost to the mobilty is achieved with SAMs with the longest alkyl chain, namely OPA, which reaches a mobility of 10-3 cm 2 /Vs, and OTS which reaches cm 2 /Vs. This is correlates with the polymer structure difference, 126

128 that is amporphous for bare and short chain SAM and partially crystalline for the long chain SAM. The value obtained on OTS is the highest ever reported for Al 2 based TFTs and it is very close to the one measured on our DDS functionalized SiO 2 based devices (about 10-2 cm 2 /Vs, also shown in Figure 2 and in general is close to the mobility reported in the literature for high MW RR-P3HT spin cast on SiO 2. This demonstrates that by means of a proper functionalization high mobilities can be obtained also with high-κ dielectrics such as Al 2. The detailed study of the dependence of the mobility upon the accumulated charge density shows that an interplay occurs between morphological order and density of states broadening in determining the mobility. Even with long chain SAM a residual broadening of the density of states exists. These findings may contribute to the ongoing debate on the origin of the inferior performance of high-κ oxides used in conjunction with organic semiconductors in TFT structures. [1] Paper submitted to JAP. on P3HT. We elaborate a strategy to prevent Al 2 diffusion into the P3HT layer by proper engineering of the polymeric layer. SOURCE Al2O3 Polymer SiO 2 GATE (Si-n) DRAIN Figure 1. Scheme of the device used in the experiments. The Al 2 was deposited by ALD at low temperature (125 C) on a bottom contact/bottom gate transistor as shown in Figure 1. The polymeric layer was either a single layer 30 nm thick of P3HT or a stack of a P3HT and poly-vinyl phenole (PVP). The PVP layer has a low free volume and is carrying OH groups that can promptly react with trimetyl aluminum (TMA) used as a precursor in the Al 2 growth process. A TEM cross section of the Al 2 /P3HT and Al 2 /PVP/ P3HT stacks are shown in Figure 2. The Al 2 /P3HT stack shows a broad interface between the oxide and the polymer with oxide inclusions into the P3HT layer. 5.2 Atomic Layer Deposited Al 2 as a capping layer for Polymer based transistors S. Ferrari 1, F. Perissinotti 1, E. Peron 1, L. Fumagalli 2, D. Natali 2 and M. Sampietro 2 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Milan, Italy 12 Politecnico di Milano, Dipartimento di Elettronica e Informazione, Milan, Italy Due to the sensitivity of organic active materials to water and oxygen, organic based Thin Films Transistors (OTFT) devices require the developement of efficient encapsulating layers. Among different passivating layers, Al 2 shows excellent gas permeation properties. we investigated the use of Atomic Layer Deposited Al 2 thin films as capping layer for OTFT. We studied in particular the mechanism of ALD Al 2 growth on the polymeric semiconductor Poly- 3 hexyltiophene (P3HT), and the effect of the growth process on the electrical properties of transistors based Figure 2. TEM cross section of the Al 2 /P3HT (left) and Al 2 /PVP/P3HT (right). 127

129 The transition between the two layers is extended over approximately 10 nm. The stack Al 2 /PVP/P3HT TEM cross section shows a well defined Al 2 /PVP interface, with a sharp contrast between the oxide and the polymer regions. In addition the interfacial roughness is extremely limited, with no indication of those Al 2 inclusions that were apparent in the Al 2 /P3HT stack. Sample μ(cm 2 V 1 s 1 ) Ion/Ioff Uncapped 1x Capped with Al x Capped with Al 2 /PVP 1.1x Table I. Electrical properties of the transistors capped and uncapped. The effect of the Al 2 layer on the electrical properties of OTFT having the structure shown in Figure 3 was studied by comparing devices uncapped, capped with Al 2 and capped with AlO/PVP. The mobilities of the transistors measured in vacuum are reported in Table I. The uncapped device has a hole mobility of 10-2 cm 2 / Vs and an on/off ratio of about 10 6, typical values for spin cast high MW P3HT. When the device is capped by Al 2, a degradation in the mobility of almost an order of magnitude occurs (Table I). This could be correlated with the diffusion of the Al 2 into the P3HT layer that might introduce trapping states into the semiconductor and increase the structural disorder. The on/off ratio is lowered by almost four orders of magnitude. The drop of the on/ off ratio can be related to the high doping caused by the Al 2 clusters embedded in the P3HT layer. Such clusters are carrying OH groups and physisorbed HO molecules that are trapped in the layer during the ALD growth process. Such species can act as dopant, increasing the off current of the transistor. Finally we tested the capping layer properties of Al 2 on our TFT devices. We measured the transfer characteristic curves of transistors in ambient air with the ones in vacuum. We can observe that in air a high doping of the uncapped device occurs through moisture and oxygen adsorption, which results in an intolerable low on/off ratio of about 3, as shown in Figure 3. On the contrary, the Al 2 /PVP capped transistor is exceptionally robust to air exposure as its transfer characteristic curves measured in vacuum and in air almost superimpose as shown in Figure 3. In addition, this device was measured after 30 days of continuous air exposure: apart from a slight degradation of the on current, we measured a negligible variation in the off current (from 43 μa to 57 μa). This confirms the exceptional properties of the ALD deposited of thin AlO films as oxygen and moisture barrier as demonstrated earlier. [1] S. Ferrari, F. Perissinotti, E. Peron, L. Fumagalli, D. Natali, M. Sampietro, Atomic Layer Deposited Al 2 as a capping layer for polymer based transistors, Organic Electronics, (2007), in press. doi: /j.orgel A novel method for the production of top contact thin film transistors based on organic semiconductors S. Ferrari 1, F. Perissinotti 1, E. Peron 1, L. Fumagalli 2, D. Natali 2 and M. Sampietro 2 1 CNR-INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy 12 Politecnico di Milano, Dipartimento di Elettronica e Informazione, Milan, Italy Figure 3. Transistor transfer characteristics of the uncapped device and the Al 2 /PVP/P3HT capped one. There is a strong demand for the development of novel methods for the fabrication of organic thin film transistors (OTFT). Since the areas of application of the 128

130 OTFT are in the low end, the methods shall be low cost, high throughput, but yet they shall meet very strict requirements in terms of minimum feature size achievable. Indeed a scalability up to 1 μm or more is desirable for practical, state of the art application. The realization of OTFT using photolithography directly on the organic semiconductor is problematic, due to the incompatibility of the solvents used during photoresist deposition/removal. We have developed a method enabling the use of conventional photolithographic processes on organic semiconducting layers. The method allows to build top gate - top contact transistors (as well as topgate bottom-contact ), enabling circuits to be designed using the same layout as in the silicon technology. In addition the gate is constituted by a thin high-κ oxide layer, therefore allowing low voltage transistor to be manufactured. The method is based on the idea of coating the Organic Semiconductor (OS) with an oxide such as Al 2 by Atomic Layer Deposition (ALD). The ALD process should be carried out to a temperature below the melting point of the OS. As compared to other protecting strategies, such as the use of thick organic insulating layers to protect the semiconductor, this method has three important advantages: 1. The etching of the oxide can be performed using aqueous solutions, that do not interact with the OS. 2. The ALD deposited oxide is pin-hole free, compact and shows excellent barrier properties against gasses and liquids. 3. The oxide itself can act as gate dielectrics with an high-κ and excellent electrical properties. Figure 2. SEM top view (left) and cross section (right) of a top contact transistor fabricated with the photolithographic method. Figure 3. Transcharacteristic of the top contact transistor. Figure 1. Scheme of the device realized with the novel method. The basic structure of the transistor that has been realized with the novel process is schematized in Figure 1. The OS chosen to study the process was Poly-3-hexyltiophene (P3HT). The gate is made by the substrate that is an n+ doped silicon wafer with 100 nm thermally grown SiO 2. The top view show that the metal contact can be defined very well on the polymer layer. The cross section shows the Au layer adhered to the P3HT layer. The transfer characteristics of the transistor are shown in Figure 3. The mobility was estimated around 10-3 cm 2 V -1 s -1, that is one order of magnitude lower than conventional bottom contact transistors. We explain this difference by considering that the fabrication process causes a slight degradation of the OS. We are currently working to improve the fabrication process to achieve the same mobility of the bottom contact transistors. 129

131 6. Very high-κ oxides for Neuroelectronics 6.1 Electrical properties of neuroelectronic devices with various interlayers between TiO 2 and p-si(100) 6.2 Energy band alignment at the TiO 2 /Si interface with various interlayers The understanding of systems composed of neuron cells, which can be contacted in semiconductor-based chips, is the basis for any development of neuronprostheses and neuron-computers. The signal transfer at the interface between semiconductor and neuron cells might occur in both directions. In this perspective, the systems investigated within the collaboration between MDM and the Department of Membrane and Neurophysics, Max Planck Institute for Biochemistry, Martinsried (Germany) MPI, consider signal transfer taking place just from semiconductor to neuron cells. A so-called action potential can be activated through an extracellular stimulation from the semiconductor. The action potential originates from electrical activity taking place in the cells as a consequence of the extracellular stimulation. The electrical activity, in turn, follows as a consequence of the opening of ion (e.g. Na, K) channels. In order to elicit action potentials in nerve cells adhered on electrodes, a certain current is required across the electrode themselves. Electrochemical reactions may cause damage to cells and electrodes. This problem is avoided by using p-si(100) as electrode. Thus, for neuroelectronic devices, p-si(100)- based electrodes are separated by insulating dielectric materials forming capacitors. Only capacitive current flows through these capacitor, and electrochemical reactions do not take place. The aim of the collaboration between MDM and MPI in this context is to 130

132 fabricate novel stimulation chips exhibiting extremely high capacitance in order promote advanced applications in neuroelectronics. These chips are first characterized, and then used for the stimulation of cells. The formerly used SiO 2 -based dielectrics are replaced by metal oxides with higher dielectric constant (κ) than SiO 2. TiO 2 was selected because of its κ value around 35 for the anatase phase, and around 80 for the rutile phase. TiO 2 films are deposited on p-si(100) using atomic layer deposition (ALD). The chips are characterized in the electrolyte/ insulator/semiconductor (EIS) capacitor configuration. Since only limited leakage and short-circuit currents play a role in EIS capacitors, the electrical characterization of high-κ TiO 2 layers is possible in more detail in EIS then in more classical metal/insulator/ semiconductor (MIS) capacitors. The bias voltage-dependent capacitances of TiO 2 films in EIS capacitors exhibit interesting properties which can be understood only with the aid of numerical calculations assuming free electrons in the TiO 2 layer. For certain bias voltage values, the low-lying conduction band of TiO 2 causes electron accumulation within the TiO 2 layer. This phenomenon seems to be at the origin of the high capacitance in TiO 2 -based EIS capacitors. These two stimulation mechanisms are ion channel opening, and electroporation of cell membranes. Depending on the stimulus shape, one of these mechanisms, or both of them simultaneously, can be induced. Since only capacitive current is flowing during stimulation, damage due to electrochemical reactions is avoided. The novel TiO 2 -based stimulation devices feature capacitances 5 orders of magnitude higher than those of SiO 2 -based ones, and are successfully used to induce two fundamental stimulation mechanisms in HEK293 cells expressing the recombinant K channel. 131

133 6.1 Electrical properties of neuroelectronic devices with various interlayers between TiO 2 and p-si(100) F. Wallrapp, 1 G. Scarel, 2 M. Perego, 2 G. Seguini, 2 M. Fanciulli, 2 and P. Fromherz 1 1 Department of Membrane and Neurophysics, Max Planck Institute for Biochemistry, Martinsried, Germany 12 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy This paper examines experimentally and with the aid of numerical simulations the electrical properties of EIS capacitors fabricated on p-si(100) and with thin TiO 2 films grown by ALD as insulating layers. The TiO 2 films were grown at 295 C in a F-120 ASM-Microchemistry Ltd. reactor on p + type ( Ωcm) Si(100) wafers using Ti isopropoxide (Ti[OCH(CH 3 ) 2 ] 4 ) as Ti precursor, and H 2 O as oxygen source. Ti[OCH(CH 3 ) 2 ] 4 was evaporated from an open boat at 40 C, while H 2 O was kept in an external reservoir at 18 C. N 2 was used both as carrier and purge gas [1]. In order to increase the total capacitance of TiO 2 -based EIS capacitors (Figure 1), various interlayers between TiO 2 and p-si(100) are considered: Si 3 N 4 [2] Al 2, and Ti-rich SiO x layers. Interlayers should have the highest possible dielectric constant in order to limit the κ value decrease of the whole dielectric stack and to reduce the equivalent oxide thickness (EOT) [3]. The Si 3 N 4 interlayer was chosen [2] because nitridated Si interfaces have a κ value of 7, higher than the one of SiO 2 [3]. The Al 2 interlayer was chosen because of its κ value of 9 [3]. The thin Si 3 N 4 and Al 2 layers were deposited, respectively, by reactive sputtering and by ALD. On the other hand, the Ti-rich SiO x layers formed unavoidably on H-terminated p-si(100) during TiO 2 layer growth. The TiO 2 κ value in the TiO 2 /Si 3 N 4 /p- Si(100) system is 31 [2]. In all TiO 2 /interlayer/p-si(100) systems (Figure 2), in the accumulation region of p- Si(100), we find saturation of capacitance. In the inversion region of p-si(100), the capacitance increases in two steps, far beyond the saturation capacitance. We explain this effect as due to electrons in the TiO 2 layer at the interface with p-si(100), whose behavior is controlled by bias voltage, by fixed positive charges at the TiO 2 /p-si(100) interface, and by alternating current (AC) voltage frequency. However, depending upon the chosen interlayer, a different band lineup between TiO 2 and p-si(100) is established [4]. Band lineup affects only slightly the saturation value of the capacitance, but greatly influences the course of the capacitance in the depletion region of p-si(100). In particular, in the TiO 2 /Al 2 /p-si(100) system, no capacitance reduction is observed at low AC voltage frequencies in the bias voltage range in which p-si(100) is in depletion (Figure 2, top). This behavior differs from the one of the TiO 2 / Si 3 N 4 /p-si(100) system, and in part also from the one of the TiO 2 /Ti-rich SiO x /p-si(100) system. On the other hand, in systems with a large number of positive fixed charges, at high AC frequencies, and at positive bias voltages, the capacitance rises steeply in the TiO 2 /Al 2 /p-si(100) and TiO 2 /Ti-rich SiO x /p-si(100) systems, not in the TiO 2 /Si 3 N 4 /p-si(100) one (Figure 2, bottom). Different distribution of interface states exchanging charges with TiO 2, with respect to the TiO 2 conduction band edge, are assumed to explain these phenomena. Our results indicate that changing the chemical nature of the interlayer at the TiO 2 /p-si(100) interface allows to control and engineer both band alignment and electrical properties of EIS capacitors of interest in neuroelectronics. Figure 1. Schematic cross section of an electrolyte/tio 2 /p-si(100) system (not to scale). The substrate is a highly p-doped Si(100) wafer. The capacitor, 2 mm in diameter, is etched in a passivating oxide of SiO 2 with a thickness of 1 μm. The TiO 2 films are fabricated by ALD on top the p-si(100) with or without interlayer. The electrolyte (160 mm NaCl) is connected with an Ag/AgCl electrode. The back of the chip is contacted by evaporated Al. DC bias voltages V ES are applied between the electrolyte and p-si(100). 132

134 [1] M. Vehkamäki, T. Hänninen, M. Ritala, M. Leskelä, T. Sajavaara, E. Rauhala, and J. Keinonen, Chem. Vap. Dep. 7, 75 (2001). [2] F. Wallrapp and P. Fromherz, J. Appl. Phys. 99, (2006). [3] G. D. Wilk, R. M. Wallace, and J. M. Anthony, J. Appl. Phys. 89, 5243 (2001). [4] M. Perego, G. Seguini, G. Scarel, F. Wallrapp, P. Fromherz, and M. Fanciulli, to be submitted. 6.2 Energy band alignment at the TiO 2 /Si interface with various interlayers M. Perego, 1 G. Seguini, 1 G. Scarel, 1 M. Fanciulli, 1 F. Wallrapp, 2 and P. Fromherz 2 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy 12 Max Planck Institute for Biochemistry, Department of Membrane and Neurophysics, Martinsried, Germany Figure 2. Area-specific differential capacitance c s of various EIS capacitors versus bias voltage V ES between electrolyte and p-si(100) at AC voltage frequency of 100 Hz. Respectively curves corresponding to first-hand (top), and conditioned chips (bottom) are shown. The type of EIS system is indicated in the legend. From the point of view of the application of chips based on the TiO 2 /interlayer/p-si(100) systems described above, it is noticeable that capacitors based on TiO 2 / Si 3 N 4 /p-si(100) systems were already practically used. Compared to them, capacitors based on the TiO 2 /Ti-rich SiO x /p-si(100) system are less suitable for advanced applications because they feature low capacitances, and also because they have a small usable bias voltage range. However, capacitors based on the TiO 2 /Al 2 /p- Si(100) system are more promising than those based on the TiO 2 /Si 3 N 4 /p-si(100) one because of the highfeatured capacitance, and the broad usable voltage range. Another important advantage of this type of capacitors is the rather smooth voltage dependent capacitance at low AC voltage frequencies, which leads to a rather constant current across the EIS when voltage ramps are applied in ranges suitable for advanced neuroelectronic applications. Few data are available in the literature on the band structure of the TiO 2 /Si heterojunction. Valence band offset (VBO) values ranging between 1 and 2.5 ev have been reported by various authors.[1-3] Variations of the VBO have been observed by Fulton et al. after post deposition annealing. [1] Band gap measurements for TiO 2 thin films and bulk TiO 2 samples have been performed by many authors. Depending on the crystallographic structure the TiO 2 films exhibit different band gap values ranging from ~ 3 ev for rutile TiO 2 to ~ 3.4 ev for anatase TiO 2, to ~ 4.0 ev for amorphous oxide. [3] No data are available on the conduction band offset (CBO) at the TiO 2 /Si heterojunction. In this work we studied systematically the band alignment at the TiO 2 /Si heterojunction combining different analytical techniques. Anatase TiO 2 films were grown on Si (100) by atomic layer deposition (ALD). Various interlayers (IL) (Si 3 N 4, Al 2 and Ti rich SiO x ) between the TiO 2 film and p-type Si (100) have been considered, and the effect of the IL on the TiO 2 /Si band line-up has been investigated. Using an electrolyte oxide semiconductor (EOS) configuration we characterized the band structure at the TiO 2 /Si heterojunction by internal photoemission spectroscopy (IPE). Two different barriers (Φ 1 and Φ 2 ) 133

135 Figure 1. Square root of the IPE yield as function of the photon energy for the for TiO 2 /Si heterojunctions with different interlayer: Ti rich SiO x, Si 3 N 4, and Al 2 at the same negative applied voltage (V) of -2.6 V. have been identified. As clearly shown in Figure 1, the Φ 1 barrier is exactly the same for all the samples while the values for the Φ 2 barrier are slightly different for samples with different IL. Optical absorption (OA) measurements are reported in Figure 2. The indirect band gap measured at the crossing point between the threshold linear fit and the background linear fit is 3.4 ± 0.1 ev. This result is compatible with data reported in the literature for TiO 2 film in anatase phase. The energy value for the Φ 1 barrier is 3.3 ± 0.1 ev, in good agreement with the band gap value obtained by OA measurements. Therefore we can assume that barrier Φ 1 corresponds to a transition from TiO 2 valence We use X-ray photoelectron spectroscopy (XPS) to determine the VBO for the TiO 2 /Si heterojunctions with the various IL. According to Kraut s method [4], a VBO value of 2.56 ± 0.05 ev was determined at the TiO 2 /Si heterojunction with Ti rich SiO x IL. Similarly we obtained a VBO value of 2.44 ± 0.05 ev and 2.73 ± 0.05 ev for the TiO 2 films grown on the Si 3 N 4 and Al 2 IL respectively. Combining XPS and IPE data we can easily calculate the CBO for the TiO 2 /Si heterojunctions with the various IL. The results of this procedure are reported in Figure 3 and clearly indicate a type II band alignment between the silicon substrate and the TiO 2 film. The CBO values at the TiO 2 /Si interface are found to be -0.2 ± 0.1 ev, -0.4 ± 0.1 ev and -0.5 ± 0.1 ev for the TiO 2 films grown on Si 3 N 4, on Ti rich SiO x and on Al 2 respectively. According to the band structure reported in Figure 3, the energy difference between the TiO 2 valence band edge and the Si conduction band edge is ~ 3.7 ev. Looking back 0.6 ev 0.4 ev TiO ev IL Si Si 3 N 4 Ti rich SiO Al 2 x CBM Si CBM ox VBM Si 3.3 ± 0.1 ev 2.73 ± 0.1 ev 2.56 ± 0.1 ev 2.44± 0.1 ev VBM ox Figure 2. Optical absorption coefficient for TiO 2 grown on quartz. Data have been fitted using an indirect gap optical model. Figure 3. TiO 2 /Si heterojunction band structure reconstructed combining XPS, IPE and OA data. A variation of the band line-up has been observed for heterojunctions with different IL. band to TiO 2 conduction band. The identification of the Φ 2 barrier is not straightforward and requires more information on the band structure at the TiO 2 /Si heterojunction. to IPE data we observe that the value for the Φ 2 barrier is approximately in the same energy range. Although the accuracy in determining the value of the Φ 2 barrier is quite limited, we can conclude that the Φ 2 threshold 134

136 is related to an electron transition from TiO 2 valence band to Si conduction band. In conclusion the combination of different analytical techniques allowed to reconstruct a coherent picture of the band structure at the TiO 2 /Si heterojunctions. Moreover our results clearly show that the band structure at a TiO 2 /Si heterojunction can be manipulated upon modification of the IL between the oxide and the underlying Si substrate. [1] C. C. Fulton, G. Lucovsky, R. J. Nemanich, Appl. Phys. Lett. 84(4), 580 (2004). [2] S. A. Campbell, D. C. Gilmer, X. Wang, M. T. Hsich, H. S. Kim, W. L. Gladfelter, and J. H. Yan, IEEE Trans. Electron Devices 44, 104 (1997). [3] C. C. Fulton, G. Lucovsky, R. J. Nemanich, J. Vac. Sci. Technol. B 20, 1726 (2002). [4] E. Kraut, R. Grant, J. Waldrop, and S. Kowalczyk, Phys. Rev. Lett. 44, 1620 (1980). 135

137 7. Metamaterials 7.1 Theory of propagation in negative refractive index media 7.2 Experiments on transmission of negative refractive index media based on Split Ring Resonators Ametamaterial is a composite medium which gains its properties from its structure rather than directly from its composition. Each cell is generally realized by dielectric and conductive elements. For wavelengths much longer than the lattice spacing, the response is averaged on the cells and the effective ε(ω), μ(ω) and n(ω) may realize values not obtainable in nature. Such materials can be designed to behave at a predictable frequency with peculiar characteristics, as negative effective permeability and permittivity materials, so the consequences for the wave propagation, predicted in 1968 by V. G. Veselago, are dramatically different with respect to ones in ordinary media. Veselago studied the electromagnetic properties of materials whose permittivity and permeability simultaneously have negative real parts and showed that these media exhibit unusual properties, such as necessary frequency dispersion of the constitutive parameters, reversal of the Doppler and Vavilov- Cerenkov effect, reversal of the Snell law, the subsequent refraction at the interface between a conventional and a negative medium, and the interchange of convergence and divergence effects in convex and concave lenses, respectively, when the lens is made left-handed. In the period , J. Pendry at Imperial College discovered how to separately reverse the electric permittivity and the magnetic 136

138 permeability by applying a new concept, based on the idea of periodic structures with highly anisotropic macroscopic cells. Such cells are able to strongly modify the constitutive constants of the materials used for the realization when the frequency of the electromagnetic wave obeys to the rule stated above between the wavelength and the cell spacing. The main property that is specific to media with negative constitutive constants is that these media support plane backward waves. For a homogeneous plane wave in a lossless material e jωt-kr with a real vector k in an isotropic positive medium, the Poynting vector S has the same direction of k. Instead, if the parameters are negative, S has the opposite direction. In 2000, J. Smith at UCLA demonstrated the reversal of both the constitutive constants at the same microwave frequency in a periodic 3-dimensional structure, obtaining negative refraction. After that experiment, several research directions have been explored. They can be divided into three major areas: negative refractive index lenses and hybrid positive/negative media for microwave propagation; radiated wave applications (antennas etc); and terahertz nanolayers and nanocircuits for optoelectronics. The constitutive constants of metamaterials are not necessarily negative, and several applications are exclusively based on the periodicity of the medium designed with suitable properties in a specific range of frequencies (quasi-zero dielectric constant, frequency selective surfaces, artificial magnetic conductors). At MDM the main interest has been devoted to guided propagation in hybrid media, from both a theoretical, numerical and experimental point of view. The features of hybrid propagation have been theoretical studied for round waveguides in the GHz domain. A special application of split ring resonators has been solved, realized and tested in the range 5-8 GHz. According to the mission of the Laboratory, innovative materials have been investigated in order to realize samples on non-conventional substrates with low losses in the Q-band frequency range. The overlap between the ability to realize samples, to predict their behavior by means of finite-elements methods, and to characterize on several frequency bands, puts the Laboratory at the state of the art of the field in those applications relevant for telecommunications and defense. Future applications on innovative resonators for paramagnetic resonance and high frequency precise dielectrometry are also expected. 137

139 7.1 Theory of propagation in negative refractive index media E. Prati 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy Part of the study on negative refraction has been devoted to propagation of modes in waveguides partially or totally filled with negative refractive index media [1]. The characteristic equation [2, 3] of a cylindrical core/cladding structure shielded by a metal wall has been derived when at least one of the two media is double negative [4]. Consider the propagation of electromagnetic waves in an isotropic cylinder of infinite length and radius R* with constants ε 1 (ω) and μ 1 (ω) immersed in a medium with ε 2 (ω) and μ 2 (ω), the latter shielded by a metal wall of radius R. Because of the translational invariance along the axis z, the fields are expressed with a dependency on z of the form e ihz, and the propagation is realized by real values of the complex wave number h. In cylindrical coordinates, the most general solution for the field E and H may be written as a linear combination of elementary functions: = ( t +Ψ z n) e ihz e il e iωt l where can be E or H, ω / 2π is the frequency of the e.m. field, l the azimuthal index, and the azimuthal coordinate. In the following, we will consider = A Z ( z nlp l βr) where Z n l is an appropriate linear combination of the J l and Y l Bessel functions, r is the radial coordinate and β 2 = ω 2 εμ h2. Once the azimuthal index l is fixed, if we want to match solutions with the same azimuthal periodicity in different materials, the selection of a discrete set of propagation modes is forced by the boundary conditions on the tangential surface separating the core from the cladding [5]. We consider also the constraint on the shield that is assumed to be a perfect metal. The continuity constraint is imposed to E z, H z, 2 E and H. We define β i = ω 2 ε i μ i h 2 where i takes on the values 1 and 2, the core/cladding interface radius R*, and the radius R of the waveguide. The derivative of the Bessel function is defined by the convention J l (β j r) = dj l (β jr). d(β j r) Propagation is allowed if the propagation constant h is the same in the two media, so that the index may be eliminated. The second condition that makes this possible is that h appears squared in all the other terms. The standard procedure described here can be applied in a straightforward manner by virtue of the two above conditions. For simplicity it is defined in the following: M 1 (r) = J l (β 2 r) J l (β 2 R) Y l (β 2 r) Y l (β 2 R), M 2(r) = J l (β 2 r) J l (β 2 R) Y l (β 2 r) Y l β 2 R) In order for the equations imposing field continuity to be linearly independent, the determinant of the system has to be zero: J l (β 1 R*) 0 0 J l (β 1 R*) lh det R *β J (β R*) i 1ω 2 l 1 J l (β 1 R*) 1 β 1 i ε ω 1 lh J l (β 1 R*) β 1 R *β J (β R*) 2 l 1 1 M 1 (R*) 0 0 M 2 (R*) lh R *β M (R*) i 2ω 2 1 M 2 (R*) 2 β 2 = 0 i ε ω 2 lh M 1 (R*) 2 β 2 R *β 2 M (R*) 2 The characteristic equation can be derived by applying a standard procedure [2, 3] that yields: ε ω 2 1 J l β 1 R * β 1 J l β 1 R * J 1 l β 1 R * β 1 J l β 1 R * ( ) ( ) ε M 2 1 ( R *) β 2 M 1 ( R *) ( ) ( ) 2M 2 ( R* ) β 2 M 2 ( R* ) = l 2 h 2 R * 2 1 β β 2 2 (4) where the dielectric (left square bracket) and the magnetic (right square bracket) factors can be considered separately to determine the modes. The last equation is obtained by imposing the compatibility of the boundary conditions on the interface between the two media, forcing h to be the same in both of them. For each l one obtains an infinite discrete set of solutions of h. For sake of clarity we will refer to the positive refractive index materials with P and to 138

140 negative refractive index materials realized with both negative ε and μ with N. An NRI core with a positive refractive index (PRI) cladding is indicated N/P, and by the same convention all the other cases, namely P/P, N/N, and P/N. A numerical approach capable in general to search the modes in the complex domain has been used to find the unknown axial propagation constants that characterize modes into the guiding structure. The propagation constants allow a full description Dielectric Constant Dielectric Constant Dielectric Constant Frequency (GHz) Frequency (GHz) Frequency (GHz) Figure 1. Cut-off frequencies (black curves) of modes in a range of negative dielectric constant in three cases: (a) N/P, (b) P/N, (c) N/N. The logarithm of the modulus of the determinant Eq. (4) at h = 0 is plotted, to evidence the curves of interest. In this system, R* = 16 mm, and R = 20 mm, l = 20; in the case (a) ε 2 = μ 2 = 1, μ 1 = -1 and ε 1 is variable, in the case b) ε 1 = μ 1 = 1, μ 2 = -1 and ε 2 is variable, and in the case c) ε 2 = μ 2 = -1, μ 1 = -1 and ε 1 is variable. of the electromagnetic field distribution of each mode present into the guiding structure, except for the relative coefficients that weight each mode with respect to the others. The zeroes of the determinant have been found by a method that consists of scanning over propagation constant at fixed frequency or, alternatively, scanning over frequency at fixed propagation constant. The cut-off frequencies are found by considering the limit frequency where the determinant vanishes when h 0. A graphical solution can be much more informative than the mere numerical values and has been obtained by plotting on a grey scale the logarithm of the absolute value of the determinant (the latter being real instead of complex because no losses have been introduced). In this way the black lines (negative infinity) show the geometrical set of points where the determinant vanishes in correspondence of the cutoff frequencies of the modes. We can compare the existence of cut-off frequency modes in the three non trivial combinations that can be realized by choosing the core and the cladding to be either positive or negative. The structure has in this case R* = 16 mm, R = 20 mm, a high azimuthal mode (l = 20) has been chosen to magnify the effects. The physical solutions are individuated by the black curves corresponding to where propagation is permitted. Figures 1 represent the logarithm of the modulus of the determinant of structures that have different fillings. An ordinary P/P filling shows a smooth behaviour of modes. Such modes never support the propagation along the waveguide under the cut-off frequency of the lowest frequency mode. The results displayed in the Figure 1 show that a different behaviour of the cut-off frequencies occurs when the fillings are N/P or P/N. On the contrary, the case N/N is smooth and similar to the P/P case. The cut-off frequency can go reach zero (no cut-off modes) at some particular values of the dielectric constant at fixed permeability. Such properties suggest intriguing applications of NRI media in guided propagation and dielectric resonators. [1] A. Alù, N. Engheta, Proc.of ICEAA 0, September 2003, Torino, Italy, pp [2] R. A. Waldron, Theory of Guided Electromagnetic Waves, Van Nostrand Reinhold Company, London, [3] G. N. Watson, A Treatise on the Theory of the Bessel Functions, Cambridge at the University Press, [4] E. Prati, Microwave Propagation in Round Guiding Structures Based on Double Negative Metamaterials, Int. Journ. Infr. And Mill. Waves, 27, , [5] M.L. Kales, Modes in a wave guide containing ferrites, J. of Applied Physics, vol. 24, pp ,

141 7.2 Experiments on transmission of negative refractive index media based on split ring resonators C. Amabile and E. Prati 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy It is well known that the equivalent electrical permittivity of a waveguide below the cutoff frequency is negative. Hence, if a negative permeability metamaterial is used to fill the waveguide, then the index of refraction below the cutoff frequency becomes negative [1]. A way to invert the sign of the permeability is to use artificial dispersive media. The Split Ring Resonators (SRR) have been introduced by J. Pendry in 1999 [2] in order to realize such kind of media. A SRR array is characterized by a strip width (in our sample 0.5 mm), a ring gap (0.2 mm), an external square perimeter (5.2 mm) and the external perimeter of an internal square (3.8 mm), with some lattice spacing. A twofold experiment has been carried out at MDM on samples based on split ring resonators (SRR). We measured the scattering parameters of a 100 mm long waveguide loaded with arrays of split ring resonators (SRR) at different densities of inclusions. Two kinds of waveguide were used in order to demonstrate the agreement with theoretical predictions and the experiment. One experiment was carried out in a standard WR137 waveguide (Figure 1), in order to demonstrate stopband effect in correspondence of the resonance frequency of the SRRs. The second experiment was done in a special waveguide machined for our experiment in order to match the WR137 flange, with a square section of mm. The waveguide was fed by a coax-to-wr137 (5.4 GHz GHz) transition. In our samples all the SRRs of all the arrays were identical. The resonant frequency of such SRR array may be calculated by using the Pendry equation [2] which gives a resonance frequency of about 8 GHz, suitable for measurements in the typical rectangular WR137 waveguide. The SRR were etched on a 1.6 mm thick and 100 mm long FR4 substrate [3]. In order to vary the density of the inclusions in the samples we varied the number of SRRs keeping the position of the first and last SRR unchanged. We fabricated and measured samples Figure 1. A double array of split ring resonators has been measured to check the propagation in a WR137 waveguide ( GHz). The same kind of SRR has been measured also in a special square waveguide with a WR137 flange, with a much higher cut-off frequency. Figure 2. Simulation of a SRR loaded square waveguide. The time averaged power flow along the waveguide axis at SRR resonance is shown. Yellow regions have negligible power flow; red and blue zones host respectively forward and backward propagation. Figure 3. The S 21 scattering parameter of the square waveguide when it is filled with the dielectric FR4 substrate slab (red) and when it is loaded by the SRR array mounted on the same substrate (gray to black). The cut-off frequency of the waveguide partially filled by the substrate is 12.5 GHz. The SRR array causes the inversion of the magnetic permeability between 8 and 11 GHz, and allows transmission below the cut-off frequency at a negative refractive index. The transmitted power is reduced as the density of inclusions decreases. 140

142 hosting from 19 to 9 SRR. The effects on the propagation of the insertion of one array into the waveguide is shown in the Figure 3, were S 21 is displayed. The red curve corresponds to the square waveguide partially filled with the FR4 substrate; the cut-off frequency is 12.5 GHz. When an array of SRR on the top of the substrate is inserted into the waveguide, the transmission is allowed where the magnetic permeability is negative. The double inversion of the dielectric constant and of the magnetic permeability between 8 and 11 GHz is demonstrated by the propagation allowed in such range even if below the cut-off frequency. Here the propagation at three different densities is shown, to demonstrate the reduction of the transmitted power as the SRR density is decreased. Further advancements are now under study in order to realize high-frequency microstructured versions of such kind of metamaterials with tunable materials. [1] R. Marques, J. Martel, F. Mesa, F. Medina, Phys. Rev. Lett. 89, (2002). [2] J. B. Pendry, A. J. Holden, D. J. Robbins, W. J. Steward, IEEE Trans. on Microwave Theory Tech. 47, 2075 (1999). [3] C. Amabile, E. Prati, Magnetic metamaterials and their dependence on the lattice spacing, submitted. 141

143 8. Tailoring innovative devices by parameters-free simulations 8.1 Semiconductiong and high-κ oxides for ultra-scaled devices and spintronics 8.2 The magnetic map of Mn-based thin film alloys on Ni substrates 8.3 Parameter free calculation of shallow states in external field 8.4 Heterojunctions for spintronic devices In the last decade, the increasing computational power of supercomputers has allowed the simulation of materials used in solid state devices by first principles (i.e. parameter-free). These ab initio techniques are, nowadays, largely used to compute structural, electronic, magnetic, and vibrational properties of metals, semiconductors, and insulators. These stateof-the art simulations on one hand allow to analyze the microscopic mechanisms responsible for the material physical properties, on the other hand they are capable to provide reliable predictions anticipating and guiding difficult experimental activities. For these reasons we have used these methods at MDM to study materials for nanoelectronic and spintronic devices. In the last years material scientists and semiconductor industries have put considerable effort to study oxides constituted by elements of the IVA column of the periodic table: TiO 2, ZrO 2, and HfO 2, three oxides extensively studied at MDM. This effort is motivated by the hardness and the high dielectric constants (high-κ) of these compounds, in view of their nanotechnology applications. For example, high-κ is a property of paramount importance for an oxide in order to substitute SiO 2 in ultrascaled nanoelectronic devices. Within this respect one of the most promising materials is hafnium dioxide, which at ambient conditions presents several metastable structures with different calculated dielectric constants. Also titanium dioxide presents polymorphism at ambient conditions. In spite of the higher dielectric constant, TiO 2 is less suitable than HfO 2 as gate oxide due to the unfavourable band offset when interfaced with IV group semiconductors leading to the high leakage current preventing TiO 2 from being used in ultra-scaled complementary metal-oxide-semiconductor (CMOS) devices and it is intensively studied for other technological applications, e.g. photochemical solar cell and optoelectronic devices. These oxides can also be efficiently used as insulating layers in magnetic tunnel 142

144 junction for spintronic applications. In fact, the desire to combine ferromagnetic and conventional semiconductors materials into new devices for electronic industry has driven increasing interest to the study of magnetic layers suitable for spin injection in semiconductors. The main target is the design and the realization of new devices based on materials with magnetic properties which could be tailored in accordance to the needs of the electronic industry and interfaced with conventional semiconductors for which the technology is well assessed. Within the solid state community is rapidly becoming popular the word spintronic [S.A.Wolf et al., Science, vol. 294, 1488 (2001)] to denote electronic-like heterostructures where the relevant physical quantity is the spin of the carriers and its interactions with external magnetic fields rather than the charge of holes and electrons and the associated electronic properties. In this context, we investigated ZnO doped with Fe impurity, that exhibits room temperature ferromagnetism, Mn-based heterojunctions such as Mn-doped GaN/AlN that presents half-metallic properties or the interface formed by the Heusler compound and conventional III-V semiconductors, for its peculiar band alignment. Nowadays, much attention is also devoted to magnetic properties of nanostructures. Indeed, metal adsorption on metallic substrates permits one to create interfaces and surface alloys that, often, have no bulk counterpart. These artificial structures have great interest because they frequently give rise to peculiar magnetic or catalytic properties, due to both atomic arrangement and reduced dimensionality. For this reason we have done a systematic calculation of Mn atoms on ferromagnetic substrates with different orientations studying the microscopic mechanisms that determine the strong outwards buckling and the magnetic properties of surface atoms. spin. Since spins obey the laws of quantummechanics, the quantum bit (qubit) associated to the spin may be used as the building block of an innovative device: the quantum computer. To realize a quantum computer it is necessary to find efficient methods for detection, manipulation, and information storage of qubits. For these reasons a method capable to compute without any adjustable parameter electronic properties of such system would be highly desiderable. In one of the following sections we will expose the theoretical results obtained at MDM by studying P-doped silicon, one of the more promising system for the realization of a solid state based quantum computer. In this chapter we will give an account of some of the materials investigated or presently under investigation at the MDM laboratory. All calculations presented are free of any adjustable parameter. With the exception of Si:P, all other results are obtained by plane wave pseudopotentials techniques, that allow to simulate in an accurate and reliable way electronic, vibrational, thermodynamic, and magnetic properties of high-κ oxides, diluted magnetic impurities in semiconductors, and heterojunctions, in order to investigate the possibility to use these materials in future nanoelectronic and spintronic ultra-scaled devices, that are of paramount interest for the electronic industry. The nuclear or electronic spin can also be used to store information, in fact the bit - the binary unit, used for computation - can, in principle, be stored in an up (or down) 143

145 8.1 Semiconducting and high-κ oxides for ultra-scaled devices and spintronics A. Debernardi, M. Fanciulli 1 CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy Within the framework of density functional theory we have computed structural and vibrational properties of high-κ dielectric oxides HfO 2 and TiO 2. We have considered different politypes of these oxides and, after structural optimization, we have computed the high-frequency dielectric constant, the effective charges, and the zone centre phonon modes necessary to simulate by first principle the static dielectric constant that is largely influenced by the vibrational contribution of the infra-red active phonon modes [1]. Our target is to investigate what are the more promising materials and what are the more suitable phases for technological application. In particular we concentrate our effort on HfO 2, one of the most promising candidate as gate oxide in ultrascaled CMOS technology (see also chapter 2.6. of the present report). In Figure 1 we display our results for phonon dispersion relations of baddeleyite HfO 2 the monoclinic phase that is stable at ambient conditions. According to our calculation [1] the static dielectric constant of the monoclinic structure is (theoretical data from Ref.[2] are in parenthesis): κ xx =17.94 (13.1), PHONON FREQUENCY (THz) Γ X S Y Γ A Figure 1. Phonon dispersions of baddeleyite HfO 2 along some high symmetry directions. κ yy =15.76 (10.8), κ zz =12.42 (7.53), κ xz =0.98 (1.82). The experimental value, from Ref. [3], is κ =21.6, obtained by electrical measurements. Calculations of the thermodynamical properties of the baddeleyite HfO 2 can be found in Ref. [1] In 1997 Parlinski [4] and co-workers predicted by firstprinciples calculations that cubic ZrO 2 is unstable at absolute zero due to soft phonon modes. In spite of this, cubic zirconia can be stabilized by adding a small fraction of Y. In the last years theoretical studies appeared computing dielectric and vibrational properties of cubic ZrO 2, HfO 2, and TiO 2 at T=0K. Since these works predicted considerably high dielectric constant for the fluorite phase, it is therefore of interest to establish if also hafnium and titanium dioxide present a similar instability. PHONON FREQUENCY (THz) L Γ K X Figure 2. Phonon dispersions of fluorite HfO 2 along some high symmetry directions. For this reason, by density functional perturbation theory, we have computed the phonon dispersion relations of the fluorite phase of HfO 2 and TiO 2. In Figure 2 we display our results for fluorite HfO 2. From the figure, one can notice the soft phonon modes (in red) in the region of the Brillouin zone between the K and the X points. In the harmonic approximation, ω 2 are the eigenvalues of the dynamical matrix, and they are usually positive; a negative eigenvalue, that gives an imaginary frequency, corresponds to a structural instability: the soft phonon mode drives a transition of the crystal to a different stable phase. According to our results the fluorite phase of pure HfO 2 has a spontaneous phase transition into a crystal with tetragonal symmetry (that is the symmetry of the soft mode). This phonon instability is similar to that displayed in Ref. [4] where the phonon dispersion curves of cubic zirconia were computed by first principles. Since the fluorite phase of ZrO 2 can be stabilized by doping with Y, due to the similarity of Γ 144

146 PHONON FREQUENCY (THz) L Γ K X Figure 3. Phonon dispersions of fluorite TiO 2 along some high symmetry directions. Figure 4. The ZnO:Fe structure. the phonon soft mode in the two cubic compounds, our results suggest that HfO 2 can be stabilized in a similar way, e.g. by doping with Y, La or Lu.[1] In Figure 3 we display our results for the phonon dispersion relations of TiO 2 along some high symmetry directions. Unlike HfO 2, the imaginary frequencies associated with soft phonon modes are not limited to the region near the X point, but extend almost in the whole Brillouin zone. This means that the tetragonal phase driven by the X-phonon would be unstable. Since the structural instability is associated with phonon modes almost in the whole Brillouin zone, we can argue that the stabilization of the cubic phase, if possible, would involve different microscopic mechanisms with respect to cubic zirconia or hafnia. High-κ dielectric oxides can also be used for spintronic applications either as tunnel oxides or when Γ doped with magnetic atoms. Magnetic impurities in semiconductors can make the compound half metallic or they can present peculiar magnetic properties, as the increase of magnetic moment of the impurity with respect to the bulk element. This effect is particularly interesting in ZnO, a wide-gap semiconductor, that has recently attracted considerable interest because it seems to exhibit ferromagnetism at room temperature. By density functional supercell calculations we computed the total energy and the magnetic properties of wurtzite ZnO doped with diluted Fe impurities in the presence of intrinsic defects. We display in Figure 4 the relaxed geometry and the magnetization density isosurface. Looking to this figure one notices the magnetic density surrounding the iron atom (with a magnetic moment of about 3.7 μ B due to the occupancy of d level) and the (small) magnetic polarization of nearest-neighbour oxygen atoms. For all defects studied we have computed the formation energy to understand how the magnetic properties at room temperature can be influenced by the presence of defects. Further, we found that Zn-vacancies occupied the next nearest neighbour to Fe atom, thus forming the complex Fe Zn -V Zn. In contrast, the complex Fe Zn -V O is energetically unfavoured with respect to the case where Fe Zn and V O are decoupled. Both type of vacancies, coupled with diluted Fe impurities, exhibit a density of states fully spin polarized at the Fermi energy, suggesting the possibility for spintronic applications. According to our results, for Zn-poor growth condition, the room temperature population of the Fe Zn -V Zn complex provides a key to understand the magnetism of this system [5]. These results will be compared with on-line Mössbauer spectroscopy experimental data obtained implanting 57 Mn at CERN within the ISOLDE collaboration. We acknowledge the Istituto Nazionale di Fisica della Materia for have provided super-computing resources through the Iniziativa Trasversale di Calcolo Parallelo. [1] A. Debernardi, M.Fanciulli, Mat. Sci. Semicon. Proc. vol. 9, (2006). [2] Zhao and D.Vanderbilt, Phys. Rev. B, vol. 65, (2002). [3] Z.J.Wang, et al., J. Crys. Growth, vol. 281, 452 (2005). [4] K. Parlinski, et al, Phys. Rev. Lett., vol. 78, 4063 (1997). [5] A. Debernardi, M.Fanciulli, Appl. Phys. Lett., vol. 90, (2007). 145

147 8.2 The magnetic map of Mn-based thin film alloys on Ni substrates Surface: Mn Ni Subsurface: Mn Ni B.R. Malonda-Boungou 1,2,3, B. M Passi-Mabiala 1 A. Debernardi 4, S. Meza-Aguilar 5, C. Demangeat Université Marien Ngouabi, Brazzaville, Congo Centre for Atomic Molecular Physics and Quantum Optics, Douala, Cameroon The Abdus Salam International Centre for Theoretical Physics, Trieste, Italy CNR-INFM Laboratorio Nazionale MDM, Agrate Brianza, Italy Universidad Autonoma de Sinaloa, Culiacan, Sinaloa, Mexico Institut de Physique et Chimie des Matériaux de Strasbourg, Strasbourg, France [011] [011] [011] Usually Mn nanostructures are non ferromagnetic so that their use for spintronics may not be suitable in the metallic form. However, due to its intrinsic high atomic moment, Mn could be very interesting, but it has to be stabilized magnetically before it can have a determining interest in the technological world. The magnetic properties of Mn are very important for the stabilization of a new class of materials, i.e., ordered two-dimensional magnetic alloys where no ordered bulk alloys exist. The case of surface alloy is very interesting, because Mn displays a high magnetic moment as well as an outward relaxation that reduces the atomic coordination, and thus enhances the magnetic moment. As a result, there is a gain in the magnetic energy which stabilizes these structures. We investigated the magnetic map of Mn thin films and MnNi ordered surface alloys on thick Ni film epitaxially grown on Cu substrates with three different crystallographic orientations: (001), (110), and (111). This is motivated by the fact that there are now considerable experimental results of Mn films on Ni substrates. LEED structure determination of compositionally ordered MnNi films epitaxially grown on Ni(001) by deposition of 3-4 Mn monolayers above 550K were performed by Wuttig et al. [1] showing that the resulting film structure is indicative of the formation of tetragonal MnNi films. A pronounced corrugation of 0.30 ± 0.02 Å was observed at the MnNi film surface (see Figure 1). Our ab initio simulations are performed within the Figure 1. Structural model of MnNi bulk alloy. (a) Overview, where the unit cell is set by the square in doted lines. (b) Interplanar distances obtained by Wuttig et al [1]. density functional theory in the generalized gradient approximation by using plane-wave pseudopotentials techniques with slab geometry and relaxing the atoms to obtain the equilibrium configuration. For the Mn monolayer on Ni(001), with two nonequivalent atoms per layer, a ferromagnetic coupling is obtained, whereas an antiferromagnetic coupling between Mn atoms is presented for Mn on Ni(011) and Ni(111). For the ordered Mn 0.5 Ni 0.5 monolayer on Ni a ferromagnetic configuration is found. At variance, for the ordered Mn 0.5 Ni 0.5 two layers thick on Ni an antiferromagnetic coupling between Mn atom nearest neighboring positions are depicted for the (001) crystallographic face whereas it is ferromagnetic for (111) and (011). Strong outwards buckling is always obtained for Mn atoms. In the following, we illustrate our data presenting results for the Mn monolayer on Ni/Cu(001) case. The Ni(001) clean surface presents a contraction of the first surface plane with 1.64 Å for the inter-layer distance between the surface and the subsurface. For a Mn mono-layer on Ni(001) a much more complex result appears (notice that there are two non-equivalent atoms per layer): half of the Mn atoms noted Mn Sa are located at 1.66 Å from the Ni subsurface layer. 146

148 The other half of Mn Sb atoms are strongly pushed outwards: they are located at 0.12 Å from the Mn Sa atoms (i.e. the corrugation of the surface is 0.12 Å). B.R. Malonda-Boungou would like to address thanks to the AIEA/ ICTP STEP-programme for financing this work at the Abdus Salam ICTP. The complete magnetic map and the distances between the various Mn and Ni layers are reported in Figure 2. In this Figure three facts are clearly shown: i) the Mn monolayer on Ni(001) presents a clear ferromagnetic behavior; ii) the magnetic moments on the Ni atoms in the subsurface layer are clearly depressed and iii) the magnetic moments on the two types of Mn atoms are dramatically different (0.58 μ B in one case, 3.38 μ B for the other). It must be remembered here that bulk Mn is of antiferromagnetic type [3] and thus, in principle, a Mn monolayer should present a purely antiferromagnetic behavior. On the other hand, Mn is in contact with Ni which is a strong ferromagnet: thus a strong induced polarization of Mn by the Nisubstrate is present. These two effects (antiferromagnetic bulk Mn and Ni induced ferromagnetism) compete and lead to the behavior displayed in Figure 2. [1] M. Wuttig, C. C. Knight, Phys. Rev. B vol. 48, (1993). [2] B.R. Malonda-Boungou, B. M Passi-Mabiala, A. Debernardi, S. Meza-Aguilar, C. Demangeat, The magnetic map of Mn-based thin film alloys on Ni/Cu(001), submitted for publication. [3] J. Hafner and D. Spisak, Phys. Rev. B, vol. 72, (2005). 8.3 Parameter free calculation of shallow states in external field A. Debernardi 1, M. Fanciulli 1, A. Baldereschi 2,3,4 Mn1 Mn2 Ni1 Ni2 [001] S S 1 S 2 S 3 S 4 S 5 S Figure 2. Magnetic moments (in μ B ) and equilibrium geometry of a Mn over-layer on Ni/Cu(001). There are two non-equivalent atoms on each layer. S-6 is the center of Ni film with 11 planes. S is the surface plane containing the Mn atoms. Magnetic moments of each atom are in circles and squares, in dashed and solid lines. Interplane distances are given in Å. However, the one (two) mono-layer(s) NiMn alloy on Ni shows a pronounced corrugation of Mn atoms in the first plane with the values of 0.26 Å (0.14 Å), in agreement with experimental results done by Wuttig et al. [1]. We believe that these transformed structures should be closely related with the magnetic properties of the film surface. We acknowledge the Consorzio interuniversitario per le Applicazioni di Supercalcolo Per Università e Ricerca (CASPUR) for computational resources provided under the project Electronic and magnetic properties on Mn-based materials and nanostructures CNR-INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy INFM-Democritos National Simulation Center, Trieste, Italy Dipartimento Fisica Teorica, Trieste University, Italy Institute of Theoretical Physics (EPFL) Lausanne, Switzerlan In 1998 B. E. Kane envisaged a scheme of a quantumcomputer [1] (i.e. a computer obeying to laws of quantum mechanics) in which the nuclear spin of 31 P impurity atom at substitutional sites of crystal silicon is used as quantum-bit (qubit). The interaction between the nuclear spins at different impurities sites is mediated by the electronic orbital of the impurity that - for shallow donors as P impurity in silicon (Si: P) - extend for few nanometers apart from the guest atom. Within this scheme, quantum bit manipulation depends of the capability of tuning by an electric field the hyperfine splitting (proportional to the square modulus of the impurity wave-function evaluated at the P nucleus) that governs the interaction between the nuclear spin of P atom and the electron spin of the hydrogen-like impurity state. The Kane s proposal [1] has driven increasing interest into the field of shallow impurities, and a number of theoretical studies that assay to compute the properties of such system were published in the last years. While shallow states confined in clusters and nanostructures of small size (i.e. up to about one nanometer) can be conveniently simulated by first principles [2], the study of larger 147

149 systems requires the use of the envelope function approximation that permits to compute, with a reduce computational effort, shallow states that extend for several nanometers. To simulate in a realistic way the ground state of P impurity in Si, it is necessary to include in the calculation the following quantities: i) band structure effects of the host material, i.e. the band anisotropy of silicon near the conduction band minima (CBM), ii) the valley-orbit interaction (VO), i.e. the coupling by the impurity potential of electronic states belonging to different degenerate CBM (valleys) of Si; iii) the central cell corrections, i.e. the difference (expected to be significant only in the central cell containing the impurity) between the true potential of the impurity and the screened Coulomb potential that is used to approximate the impurity potential. In the past, within the envelope function approximation, different attempts appeared in literature to compute P-impurity states in bulk Si, the majority neglects the valley orbit-interaction, and few take into account it in an approximate or phenomenological way [3]. So far, a reliable approach including both the valley-orbit interaction and the central cell correction to compute electronic properties of shallow impurities in external electric field is missing. By using a Gaussian basis set we have developed a robust and efficient method, based on the envelope function of the conduction band, to compute shallow impurity states in semiconductors without adjustable parameters. In our approach we take into account the band anisotropy within the effective mass approximation, and we are able to compute exactly, within the numerical accuracy due to the use of a finite basis set, the valley-orbit interaction of a realistically screened coulomb potential, of the core potential, and of the electric field. We have computed the energy levels of the shallow P impurity in silicon crystal, the hyperfine splitting of the ground state and their dependence on the uniform electric field applied along the [001] direction. We have shown that in order to reproduce correctly the Si:P ground state one has to include the central cell corrections contribution due to impurity core electrons, a quantity that, to the best of our knowledge, has not been taken into account before in this type of calculations. We found that the ground state energy decreases by increasing the magnitude of the field, and that at high electric field (~ 2-3 MV/m) the spectrum narrowing of 1s manifold, predicted in Ref. [3], is quite small, while the main effect is the mixing of s- and p-like states that, at a critical value of the field E cr =2.5 MV/m, leads to the vanishing of the hyperfine splitting. Further, we predicted the dependence of super-hyperfine splitting ENERGY (mev) Ψ(0) ELECTRIC FIELD (MV/m) ELECTRIC FIELD (MV/m) Figure 1. Top panel: solid lines denote the lower energy levels of Si: P as a function of uniform electric field directed along (001); a dashed line denotes the minimum energy of the barrier surrounding the impurity. Bottom panel: the square modulus of the ground state wavefunctions (dotted line) computed at the impurity site (and normalized to zero field result). Figure 2. Hyperfine splitting as a function of nano-particle diameter. Full diamonds denote experimental values obtained by M.Fujii and co-workers [5]. of the A-shell as a function of electric field. In Figure 1 (top panel) we display our results for the computed impurity level as a function of the applied electric filed As usual, we take the zero of the energy corresponding to the value of the conduction band minima (at zero electric field) of silicon. In the bottom panel of Figure 1 we display the square modulus of the ground state wave-functions computed at the impurity site Ψ(0) 2, i.e. a quantity proportional to the hyperfine splitting. Our approach is capable to reproduce the effect of confinement in P doped Si nanoparticles, as proven in Figure 2, where we display our results for the hyperfine splitting as a function of nanoparticle diameter and we compare them with experimental data taken from 148

150 literature. The agreement with experiment is excellent, taking into account that the theoretical results are obtained without adjustable parameters (effective masses and dielectric constant of bulk-si are the only experimental data used in our simulation). [1] B. E. Kane, Nature, vol. 393, 133 (1998). [2] D. V. Melnikov and J.R. Chelikowsky, Phys. Rev. Lett. vol. 92, (2004). [3] M. Friesen, Phys.Rev.Lett., vol. 94, (2005). [4] A. Debernardi, A. Baldereschi, and M. Fanciulli, Phys. Rev. B, vol. 74, (2006). [5] M. Fujii, A. Mimura, S. Hayashi, Y. Yamamoto and K. Murakami, Phys. Rev. Lett., vol. 89, (2002). build an interface that conserves the magnetization of bulk Heusler compound. The band alignment of these heterostructures does not present any trivial trends due to the influence of the chemical composition of the substrate. Our results can be extended to describe the band alignment of bulk NiMnSb with bulk GaAs or GaSb, while in the case of InP substrate we found that the band alignment at the junction is strongly influenced by the superlattice period. We have analyzed the microscopic mechanisms that produce the lost of the half-metallic behavior at the interfaces, and we have discussed limitations and difficulties to restore the halfmetallicity at the interface to make the heterostructure suitable to inject spin polarized into a semiconductor.[3] In a recent experiment, the magnetic properties of 8.4 Heterojunctions for spintronic devices A. Debernardi 1, M. Peressi 2,3, A. Baldereschi 2,3, CNR-INFM, Laboratorio Nazionale MDM, Agrate Brianza, Italy INFM-Democritos National Simulation Center, Trieste, Italy Dipartimento Fisica Teorica, Trieste University, Italy Institute of Theoretical Physics (EPFL) Lausanne, Switzerland Recently, NiMnSb has been successfully grown epitaxially on several substrates, such as GaAs(001) and InP(001), however, theoretical studies, based on density functional theory, have shown that, in general, the interfacial layers of the Heusler compound lost their half-metallic properties, i.e. the density of state at the Fermi energy become unpolarized, forbidding the spininjection. We have studied junctions formed by III-V semiconductors and half-metal Heusler compound to understand the microscopic mechanisms responsible of the band alignment of the heterostructure in order to design a junction suitable for spin injection [1]. By means of super-cell geometry that simulate a long period superlattice, we have computed [2,3] the structural, magnetic and electronic properties of the junctions formed by epitaxially NiMnSb on a substrate (001) oriented constituted by GaAs, InP or GaSb. We have analyzed the role of different chemical composition of the substrate and of the effect of different lattice parameters to determine the band alignment and the magnetic properties of the interface. We found the appropriate atomic termination of the junction that allows us to Figure 1. The AlN/GaN(0001) heterostructure. a digital ferromagnetic heterojunction Ga 1-x Mn x N/ GaN obtained by molecular beam epitaxy have been investigated. This system is very interesting for spintronic applications because it involves wideband-gap semiconductors. For this reason, we have investigated heterostructures based on wurtzite-type materials computing, after structural optimization, the magnetic properties of the junction AlN/GaN(0001) with different Mn concentration. We found [4] that this structure presents a completely polarized density of spin at the Fermi energy making it suitable to spin injection. In Figure 1 we present one of the structures studied. We acknowledge the Istituto Nazionale di Fisica della Materia for have provided super-computing resources through the Iniziativa Trasversale di Calcolo Parallelo. [1] A. Debernardi M. Fanciulli, First principles Calculation of Materials for Spintronics and Nanoelectronic Devices, Science and Supercomputing at CINECA, 2005 report, 453. [2] A. Debernardi, M. Peressi, and A. Baldereschi, Comp. Mater. Sci., vol. 33, 263 (2005). [3] A. Debernardi et al., to be published. [4] A. Debernardi, Superlattices Microst. vol. 40, (2006). 149

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153 ø Edited by: Marco Fanciulli Alberto Debernardi Sandro Ferrari Anna M. Ferretti Mara Lanati Enrico Prati Giovanna Scarel Sabina Spiga Grazia Tallarida Claudia Wiemer ø Photographs by: Roberto Mantovan ø IT support: Roberto Colnaghi 152