Book of abstracts - Euro-MBE 2011 - CNRS

16th
European
MBE
2011
Book of Abstracts
Photo: Laurent SALINO / Alpe d’Huez Tourisme
16 th European Molecular Beam Epitaxy Workshop
March 20 th -23 rd , 2011, Alpe d’Huez, France

16th European Molecular Beam Epitaxy Workshop
Book of
Abtracts
March 20 – 23, 2011
Congress Center, Alpe d’Huez
France

The Euro-MBE Workshop is held biennially and began 30 years ago
in Germany (1st European Workshop on MBE, Stuttgart, April 1981). It
stands as one of the most renowned and prestigious scientific meetings on
MBE. The last three Workshops were held in Grindelwald, Switzerland
(2005), Sierra Nevada, Spain (2007), Zakopane, Poland (2009) and were
highly successful, both from the participation viewpoint and from the high
quality scientific standards.
Topics
- III-V, II-VI Materials and Heterostructures
- Wide Bandgap Materials (III-Nitrides, Oxides, SiC)
- Si, SiGe and Related Materials
- Ferromagnets and Spintronics
- Novel Materials
- Low Dimensional Structures (Nanowires, Quantum Dots, …)
- Fundamentals of MBE Growth, in-situ Monitoring
- Devices and MBE production issues

Organizing Committee
Serge TATARENKO (Chairman), Régis ANDRE, Edith BELLET-AMALRIC,
Bruno DAUDIN, Véronique FAUVEL, Yann GENUIST, and Eva MONROY
Technical support: Yoann CURE, Marion DUCRUET, Jean DUSSAUD
Affiliation: CEA-CNRS-UJF group "Nanophysique et Semiconducteurs"
Institut Néel - CNRS and INAC-CEA Grenoble (France)
Program Committee
• M. BUGAJSKI, Institute of Electron Technology, Warsaw (Poland)
• T. FOXON, Univ. of Nottingham (UK)
• N. GRANDJEAN, Ecole Polytechnique Federale de Lausanne (Switzerland)
• D. GRÜTZMACHER, Institut für Bio- und Nanosysteme, Jülich (Germany)
• M. HOPKINSON, Univ. of Sheffield (UK)
• S. IVANOV, Ioffe Phys-Technical Institute, St Petersburg (Russia)
• J. MASSIES, CRHEA-CNRS, Valbonne (France)
• J. OSTEN, Leibniz University, Berlin (Germany)
• M. PESSA, Tampere Univ. of Technology (Finland)
• H. RIECHERT, Paul-Drude-Institut, Berlin (Germany)
• C. SKIERBISZEWSKI, Unipress, Warsaw (Poland)
• L. SORBA, Instituto Nanoscienze-CNR, Pisa (I) (Italy)
• G. SPRINGHOLZ, Johannes Kepler Univ., Linz (Austria)
• E. TOURNIE, Univ. Montpellier 2 (France)
• S. WANG, Chalmers Univ. of Technology, Gothenburg (Sweden)
• W. WEGSCHEIDER, ETH Zürich (Switzerland)
Contact Address
Véronique FAUVEL Tel: +33 476 88 10 89
Secrétariat Dépt. Nano Fax: +33 476 88 11 91
Institut Néel – CNRS
E-Mail : embe2011@grenoble.cnrs.fr
25 rue des Martyrs, B.P. 166 http://embe2011.neel.cnrs.fr/
38042 Grenoble, France

Invited speakers
D. AS, Paderborn Univ. (Germany)
Recent device applications of non-polar cubic group III-nitrides
O. BIERWAGEN, Univ. California Santa Barbara (USA)
MBE of semiconducting oxides
A. BONANNI, Johannes Kepler Univ., Linz (Austria)
Controlling and visualizing the distribution of transition metals in nitrides
J.-M. CHAUVEAU, CRHEA-CNRS (France)
Polar and nonpolar (Zn,Mg)O/ZnO heterostructures : the benefits of
homoepitaxy
M. EIBELHUBER, Johannes Kepler Univ., Linz (Austria)
MBE growth of IV-VI quantum dots
F. FURTMAYR, Walter Schottky Institut, Munich (Germany)
Optical and structural properties of III-Nitride nanowires and nanowire
heterostructures
F. GLAS, LPN-CNRS (France)
Growth kinetics of III-V nanowires
M. GUINA, Tampere Univ. of Technology (Finland)
Recent advances in MBE of dilute-nitrides and related device applications
V. NOVAK, Institute of Physics ASCR, Prague (Czech Rep.)
MBE growth of LiMnAs
J.-B. RODRIGUEZ, Univ. Montpellier 2 (France)
Current developments in MBE growth of highly mismatched materials
S. SANGUINETTI, Univ. degli Studi Milano-Bicocca (Italy)
GaAs based nanostructures grown by droplet epitaxy
T. WIETLER, Leibniz Univ. Hannover (Germany)
Surfactant-modified epitaxy of germanium layers on silicon for high
mobility channels

Sponsors
http://www.home.agilent.com/
http://www.omicron.de/
http://www.azeliselectronics.com/
http://www.riber.com/
http://rta-instruments.com/
http://www.dca.fi/
http://eurotherm.com/
http://www.staibinstruments.com/
http://www.hidenanalytical.com/
http://www.surface-tec.com/ /
http://www.mbe-kompo.de/
http://www.vbseurope.com/
http://www.oerlikon.com/leyboldvacuum/
http://www.veeco.com/
http://www.vinci-technologies.com/
http://www.rhonealpes.fr/
http://www.fondation-nanosciences.fr/

List of Exhibitors present on site, sponsors and the booth # :
http://www.oerlikon.com/leyboldvacuum/
#22
http://altec-equipment.com/
#20
http://www.mbe-kompo.de/
#18
http://www.laytec.de/
#12
http://www.veeco.com/
#7
http://www.mewasa.ch/
#11
http://www.vbseurope.com/
#6
http://www.axt.com/
#2
http://www.dca.fi/
#19
http://www.samtelgmbh.com/
#10
http://www.riber.com/
#14, #15
http://www.mcse.fr/
#4
http://www.omicron.de/
#17
http://www.inficon.com/
#5
http://rta-instruments.com/
#13
http://www.wafertech.co.uk
#21
http://eurotherm.com/
#8
http://www.createc.de
#1
http://www.staibinstruments.com/
#16
http://www.wepcontrol.com/
#3
http://www.cvtechnology.com
#4
http://www.vinci-technologies.com/
#9

Main locations
for Euro-MBE 2011
A : Hotel les Grandes Rousses
B : Pierre & vacances
C : Hotel Pic Blanc
: (33) 4 76 11 42 42
D : Palais des Congrés
A
D
B
C

Conference Program

Program of the 16th European Molecular Beam Epitaxy Workshop
March 20 th -23 rd , 2011, Alpe d’Huez
Sunday, March 20 th
11:00 – 20:00 Registration Lobby of Hotel Pic Blanc
18:00 – 19:15 Welcome glass of wine
Hotel Pic Blanc
19:30 – 21:30 VEECO USERS’ MEETING
Hotel Pic Blanc
Time
Monday, March 21 st
8:15-8:30 OPENING SESSION
GaAs based nanostructures grown by droplet epitaxy
Monday
8:30-9:00
Mo1.1
(invited)
S.Sanguinetti, C. Somaschini, S. Bietti and N. Koguchi
L-NESS and Dip. di Scienza dei Materiali, Università di Milano Bicocca, Italy
Monday
9:00-9:15
Mo1.2
Interfacial strains in InAs/AlSb multilayers for short
wavelength quantum cascade lasers
C. Gatel, B. Warot-Fonrose, A. Ponchet, C. Magen, R. Ibarra, R.
Teissier and A.N. Baranov
CEMES-CNRS, Toulouse, France
Monday
9:15-9:30
Mo1.3
Monday
9:30-9:45
Mo1.4
Monday
9:45-10:00
Mo1.5
Arsenides I
200 mm GaAs wafers by MBE on SGOI and Ge/Si substrates
M. Richter, T. Topuria, C. Marchiori, M. El-Kazzi, C. Rossel, C. Gerl,
D.J. Webb, T. Smets, C. Andersson, M.Sousa, D. Caimi, L.
Czornomaz, H. Siegwart, J.-F. Damlencourt, J.-M. Hartmann, P.M.
Rice, and J. Fompeyrine
IBM Research – Zurich, Switzerland
Tuning the size, strain and band offsets of InAs/GaAs
quantum dots through a thin GaAs(Sb)(N) capping layer
J.M. Ulloa, M. Montes, K. Yamamoto, A. Guzman, A. Hierro, M.
Bozkurt, P.M. Koenraad, D. Fernández, D. González and D. Sales
ISOM and Dpto. Ing. Electronica, Univ. Politecnica Madrid, 28040 Spain
InAs Quantum Dot Chains Grown on Nanoimprint Lithography
Patterned GaAs(100)
T. V. Hakkarainen, A. Schramm, J. Tommila, A. Tukiainen, R.
Ahorinta, M. Dumitrescu, M. Guina.
Optoelectronics Research Centre, Tampere University of Technology,Finland
10:00-10:30 COFFEE BREAK

MBE growth of LiMnAs
Monday
10:30-11:00
Mo2.1
(invited)
V. Novak
Institute of Physics ASCR, Prague, Czech Rep.
Monday
11:00-11:15
Mo2.2
Monday
11:15-11:30
Mo2.3
Monday
11:30-11:45
Mo2.4
Monday
11:45-12:00
Mo2.5
Monday
12:00-12:15
Mo2.6
Antimonides and Phosphides
Growth and structural properties of InSb-based
heterostructures
A.N. Semenov, B.Ya. Meltser, V.A. Solov'ev, T.A. Komissarova,
Ya.V. Teren't'ev, A.A. Sitnikova, and S.V. Ivanov
Ioffe Physical-Technical Institute of RAS, St. Petersburg , Russia
Comparison of interface properties in two-dimensional
heterostructures grown layer-by-layer or by step flow
E. Luna , R. Hey, A. Guzmán and A. Trampert
Paul-Drude Institut für Festkörperelektronik , Berlin, Germany
MBE growth of (GaAsPN/GaPN)/GaP quantum wells light
emitting diode
A. Bondi, C. Cornet, W. Guo, O. Dehaese, N. Chevalier, C. Robert,
T. Nguyen Thanh, S.Richard, M. Perrin, A. Létoublon, J. P. Burin. J.
Even, O. Durand and A. Le Corre
Université Européenne de Bretagne, INSA, FOTON, RENNES, France
Lattice mismatch accommodation at the GaSb/GaAs and
GaSb/GaP interfaces
L. Desplanque, S. El Kazzi , C. Coinon, Y. Wang, P. Ruterana and X.
Wallart
IEMN, CNRS and University of Lille , France
Defects in MBE grown InAs/GaSb superlattices
on GaSb substrates
M. Walther, R. Rehm1, J. Schmitz, J. Niemasz, F. Rutz, A. Wörl, L.
Kirste, R. Scheibner, J. Ziegler, and A. Danilewsky
Fraunhofer-Institut für Angewandte Festkörperphysik Freiburg, Germany
BREAK 12:15 – 17:00
Lunch at Hotel Pic Blanc or in a mountain restaurant of the resort + Free time

Making nitrides magnetic
Monday
17:00-17:30
Mo3.1
(invited)
A. Bonanni
Institut für Halbleiter- und Festkörperphysik, Johannes Kepler University,
Linz - Austria
Monday
17:30-18:00
Mo3.2
(invited)
Optical and structural properties of III-Nitride nanowires and
nanowire heterostructures
F. Furtmayr
Walter Schottky Institut, Munich , Germany
Monday
18:00-18:15
Mo3.3
Monday
18:15-18:30
Mo3.4
Monday
18:30-18:45
Mo3.5
Monday
18:45-19:00
Mo3.6
Nitrides
Specialized MBE system for growth of high quality III-N
heterostructures
A. Alexeev , D. Krasovitsky, S.Petrov and V. Chaly
SemiTEq JSC, St.Petersburg , Russia
Improved luminescence and thermal stability of MBE-grown
semipolar (11-22) InGaN quantum dots
A.Das, Y. Kotsar, A. Lotsari, Th. Kehagias, Ph. Komninou, and E.
Monroy
CEA-CNRS group « Nanophysique et Semiconducteurs », Institut
Néel/CNRS-Univ. J. Fourier and CEA, INAC, SP2M, Grenoble, France
Group III-nitrides growth on N-polar substrates
C. Cheze, M. Sawicka, M. Siekacz, H. Turski, G. Cywiński, B.
Grzywacz, S. Grzanka, I. Dzięcielewski, B. Łucznik,
M. Boćkowski, C. Skierbiszewski
TopGaN Ltd, Warszawa, Poland
MBE growth of GaN using in-situ SiN treatments
F. Semond, E. Frayssinet, M. Leroux, Y. Cordier, M. Réda Ramdani,
J.C. Moreno, S. Sergent, B. Damilano, P. Vennéguès, O. Tottereau
and J. Massies
CRHEA/CNRS, Sophia Antipolis, Valbonne, France
POSTER SESSION MoP1
19:00 – 20:45
Cocktail buffet: regional specialties in the exhibition room
21:00
RIBER USERS’ MEETING

Program of the 16th European Molecular Beam Epitaxy Workshop
March 20 th -23 rd , 2011, Alpe d’Huez
Time
Tuesday, March 22 nd
Advances of dilute-nitrides MBE technology and related
device applications
Tuesday
8:30-9:00
Tu1.1
(invited)
M. Guina, V.-M. Korpijärvi, J. Puustinen, A. Aho, T. Leinonen, and A.
Tukiainen
Tampere Univ. of Technology , Finland
Tuesday
9:00-9:15
Tu1.2
Current-injection lasing in GaAs quantum dots grown by
droplet epitaxy
M. Jo, T. Mano and K. Sakoda
National Institute for Materials Science, Tsukuba, Japan
Tuesday
9:15-9:30
Tu1.3
Tuesday
9:30-9:45
Tu1.4
Tuesday
9:45-10:00
Tu1.5
Arsenides II
MBE growth of high power Modelocked Integrated External-
Cavity Surface Emitting Laser (MIXSEL) with 6.4 W output
power
M. Golling, B. Rudin, V. J. Wittwer, D. J. H. C. Maas, Y. Barbarin, M.
Hoffmann, O. D. Sieber, T. Südmeyer, U. Keller
Department of Physics, Ultrafast Laser Physics Lab, ETH Zurich, Switzerland
1.55 μm lasers based on shape-engineered
InAs/InAlGaAs/InP (100) quantum dots
C. Gilfert, V. Ivanov and J.P. Reithmaier
Technische Physik, Institute of Nanostructure Technologies and Analytics,
University of Kassel, Germany
Broadband emission and mode-locking using controlled
distributions of InGaAs quantum dots
M. Hopkinson, M. Hugues , P.D.L. Greenwood , M.Krakowski, M.
Calligaro, S.Bruer, , M. Rossetti, W. Elsäßer, I. Montrosset
Department of Electronic and Electrical Engineering, University of Sheffield,
UK
10:00-10:30 COFFEE BREAK

Tuesday
10:30-11:00
Tu2.1
(invited)
Current developments in MBE growth of highly mismatched
materials
J.-B. Rodriguez
Univ., Montpellier 2, France
MBE of semiconducting oxides
Tuesday
11:00-11:30
Tu2.2
(invited)
O. BIERWAGEN
Univ. California Santa Barbara , USA
Tuesday
11:30-11:45
Tu2.3
Tuesday
11:45-12:00
Tu2.4
Tuesday
12:00-12:15
Tu2.5
New trends in MBE
MBE growth of the topological insulator Bi2Te3 on Si (111)
Substrates
G. Mussler, J. Krumrain , L. Plucinski , and D. Grützmacher
Institute of Bio- and Nanosystems 1, Research Center Jülich, Germany
Atomic-scale mapping of quantum dots using direct x-ray
methods
R. Clarke, D.P. Kumah ,+, R.S. Goldman , V. Dasika , C. Schlepütz ,
Y. Yacoby, E. Cohen , and Y. Paltiel
University of Michigan, Department of Physics, Ann Arbor, USA
II-VI-based microcavities for the blue-violet spectral range
S. Klembt, C. Kruse , M. Seyfried , K. Sebald , J. Gutowski and D.
Hommel
Institute of Solid State Physics, Semiconductor Epitaxy, University of
Bremen, Germany
BREAK 12:15 – 17:00
Lunch at Hotel Pic Blanc or in a mountain restaurant of the resort + Free time

Tuesday
17:00-17:30
Tu3.1
(invited)
Recent device applications of non-polar cubic group IIInitrides
D. As
Paderborn Univ., Germany
Tuesday
17:30-18:00
Tu3.2
(invited)
Polar and nonpolar (Zn,Mg)O/ZnO heterostructures : the
benefits of homoepitaxy
J.-M. Chauveau
CRHEA-CNRS, Sophia Antipolis, France
Tuesday
18:00-18:15
Tu3.3
Tuesday
18:15-18:30
Tu3.4
Tuesday
18:30-18:45
Tu3.5
Tuesday
18:45-19:00
Tu3.6
Wide bandgap
A bi-layer oxide buffer approach for the integration of single
crystalline GaN on Si (111) platform
L. Tarnawska, P. Zaumseil, M. Kittler, P. Storck, R. Paszkiewicz,
and T. Schroeder
IHP, Im Technologiepark 25, Frankfurt (Oder), Germany
GaN/AlGaN superlattices grown by PAMBE for intersubband
applications in the infrared spectral range
Y. Kotsar, A. Das, E. Bellet-Amalric, E. Sarigiannidou, H.
Machhadani, S. Sakr, M. Tchernycheva, F. H. Julien and E. Monroy
CEA-CNRS group « Nanophysique et Semiconducteurs », Institut
Néel/CNRS-Univ. J. Fourier and CEA, INAC, SP2M, Grenoble, France
Growth and Characterization of ZnO/ZnMgO Quantum Wells
B. Laumer, Fabian Schuster, Thomas Wassner, Martin Stutzmann,
and Martin Eickhoff
Walter Schottky Institut, Technische Universität München, Garching,
Germany
InGaN laser diodes operating at 450-455 nm grown by RF-
Plasma MBE
M Siekacz, H. Turski, M. Sawicka, G. Cywiński, J. Smalc-
Koziorowska, P. Wiśniewski, P. Perlin, I. Grzegory and C.
Skierbiszewski
Institute of High Pressure Physics, Polish Academy of Sciences, Warszawa,
Poland
POSTER SESSION TuP2
19:00 – 20:45
Cocktail buffet:regional specialties in the exhibition room
21:00
WORKSHOP BANQUET

Program of the 16th European Molecular Beam Epitaxy Workshop
March 20 th -23 rd , 2011, Alpe d’Huez
Time
Wednesday, March 23 rd
Surfactant-modified epitaxy of germanium layers on silicon
for high mobility channels
Wednesday
9:00-9:30
We1.1
(invited)
T. Wietler
Leibniz Univ. Hannover , Germany
Wednesday
9:30-9:45
We1.2
Wednesday
9:45-10:00
We1.3
Wednesday
10:00-10:15
We1.4
Group IV materials
Epitaxial growth of SrTiO3 on Si : strain relaxation and
formation of tetragonal domains
G. Saint-Girons, G. Niu, J. Penuelas, L. Largeau, B. Vilquin, J.L.
Maurice C. Botella and G. Hollinger
Université de Lyon, Institut des Nanotechnologies de Lyon, Ecole Centrale
de Lyon, France
Morphology and luminescence properties of Sb mediated
Ge/Si quantum dots
A.A. Tonkikh, N.D. Zakharov, V.G. Talalaev, A.V. Novikov, K.
Kudryavtsev,B. Fuhrmann, H.S. Leipner, P. Werner
Max-Planck Institute of Microstructure Physics, Halle, Germany
Growth of small-period Si/Ge quantum dot crystals by MBE
S. Borisova, C. Dais ,J.C. Gerharz , G. Mussler and D. Grützmacher
Institute of Bio- and Nanosystems 1, Forschungszentrum Jülich, Germany
Wednesday
10:15-10:30
We1.5
In situ STM and RHEED study of tensile strained Si grown on
Ge (001) substrates
B. Sanduijav, D. Matei, and G. Springholz
Institut für Halbleiter- und Festkörperphysik, Johannes Kepler University,
Linz, Austria
10:30-11:00 COFFEE BREAK

Wednesday
11:00-11:30
We2.1
(invited)
MBE growth of IV-VI quantum dots for MIR devices
M. Eibelhuber, A. Hochreiner, T. Schwarzl, H. Groiss, W. Heiss1, G.
Springholz ,V. Kolkovsky, G. Karczewski, and T. Wojtowicz
Johannes Kepler Univ., Linz , Austria
Wednesday
11:30-11:45
We2.2
Wednesday
11:45-12:00
We2.3
Wednesday
12:00-12:15
We2.4
Wednesday
12:15-12:30
We2.5
Wednesday
12:30-12:45
We2.6
Devices
The role of doping scheme in ultra-low disorder MBE grown
mesoscopic FQHE devices
V. Umansky , M. Heiblum, M. Dolev and Y. Gross
Braun Center for Submicron Research, Weizmann Institute of Science,
Israel
Investigations of Si-dopant layers on ultrahigh-mobility 2
DEGs in GaAs/AlGaAs-structures
C. Reichl, E. de Wiljes, C. Rössler and W. Wegscheider
ETH Zürich, Laboratorium für Festkörperphysik, 8093 Zürich, Switzerland
Short wavelength high power Quantum Cascade Lasers
X. Marcadet, M. Carras, B. Simozrag, M. Garcia,
G. M. De Naurois, G. Maisons, O. Parillaud, O. Patard, F.
Pommereau, O. Drisse, F. Alexandre, J. Massies
Alcatel Thales III-V Lab, 91767 Palaiseau cedex, France
MBE growth of InGaAs/GaAsSb based mid-infrared and THz
quantum cascade lasers
H. Detz, A.M. Andrews , P. Klang , C. Deutsch , M. Nobile , W.
Schrenk , K. Unterrainer and G. Strasser
Center for Micro- and Nanostructures and Institute for Solid-State
Electronics, Vienna University of Technology, 1040 Wien, Austria
Room temperature operation of a GaInAsSb/AlGaInAsSb
digital alloy laser diode at 3.3 μm
S. Belahsene, K. S.Gadedjisso1, G. Boissier1, P. Grech1, G. Narcy
and Y. Rouillard
Institut d'Electronique du Sud, UMR 5214 CNRS, Université Montpellier 2,
Montpellier, France
BREAK 12:45–14:45
Lunch at Hotel Pic Blanc

Growth kinetics of III-V nanowires
Wednesday
14:45-15:15
We3.1
(invited)
F. Glas
LPN-CNRS, Marcoussis, France
Wednesday
15:15-15:30
We3.2
Wednesday
15:30-15:45
We3.3
Wednesday
15:45-16:00
We3.4
Wednesday
16:00-16:15
We3.5
Wednesday
16:15-16:30
We3.6
Nanowires
InAs Quantum Dot Arrays Decorating the Facets of GaAs
Nanowires
E. Uccelli, J. Arbiol , J.R. Morante , A. Fontcuberta i Morral
Laboratoire des Matériaux Semiconducteurs, Ecole Polytechnique Fédérale
de Lausanne, Switzerland
AlAs-GaAs core-shell nanowires grown by chemical beam
epitaxy
A. Li, D. Ercolani , F. Rossi , L. Nasi , G. Salviati, F. Beltram and L.
Sorba
NEST, Istituto Nanoscienze-CNR and Scuola Normale Superiore, Pisa, Italy
Distinct nucleation and growth modes of self-assisted InAs
nanowires on bare Si(111)
E. Dimakis, J. Lähnemann, U. Jahn, S. Breuer, M. Hilse, L.
Geelhaar, and H. Riechert
Paul Drude Institute for Solid State Electronics, Berlin, Germany
Strain balanced technique for the growth of very high aspect
ratio quantum posts
D. Alonso-Álvarez, B. Alén, J. M. Ripalda, J. Llorens, A. G.
Taboada, Y. González, L. González, F. Briones, M. A. Roldán, J.
Hernandez-Saz, M. Herrera and S.I. Molina
IMM-Instituto de Microelectrónica de Madrid, Spain
Polarity of GaN nanowires grown by Plasma-Assisted
Molecular Beam Epitaxy
K. Hestroffer, C. Bougerol, C. Leclere, H. Renevier, J. L. Rouvière
and B. Daudin
CEA-CNRS group « Nanophysique et Semiconducteurs », Institut
Néel/CNRS-Univ. J. Fourier and CEA, INAC,SP2M, Grenoble, France
16:30-17:00
Award ceremony - Closing address
Free time
Free access to keep-fit center, swimming pool, spa
at Hotel Pic Blanc
Starting 19:00
Farewell Dinner

RHEED
S T A I B INSTRUMENTS
Powerful In-Situ Growth Characterization
PEEM
S T A I B INSTRUMENTS
PhotoElectron Emission Microscope
For MBE, PLD,
Laser MBE, CVD
etc.
Growth control
Oscillation detection
Latice spacing
UHV and high pressure
Visit
us
at
www.staibinstruments.com
Electron Sources
Nanotechnology
Surface Analysis
Scanning Imaging
Space Simulation
SUPER
Cylindrical Mirror Analyzer
Energy Analzyzer AES / XPS / UPS
Thermal Processing
Evaporation
Ionization

Poster sessions
Monday & Tuesday

POSTER SESSION 1
Monday, March 21 st 19:00 – 20:45
OPENING SESSION
MoP01
Post-heat treatment on the improvement of efficiency in
CdTe thin film solar cells.
Deliang Wang ,Zhizhong Bai.
MoP02
MoP03
MoP04
MoP05
Sn doped GaAs by CBE using tetramethyltin.
C. García Núñez, D. Ghita, B. J. García.
Early stages of the growth of InP and GaAs islands on
SrTiO3 substrates.
B. Gobaut, J. Penuelas, A. Chettaoui, J. Cheng, G. Grenet,
L. Largeau and G. Saint-Girons.
Growth directions and structural properties of InP
nanowires fabricated on Si and SrTiO3 substrates.
J. Penuelas, K. Naji, H. Dumont, G. Saint-Girons, G. Patriarche,
M. Gendry.
Optimization of MBE-grown AlSb/InAs High Electron
Mobility Transistor Structures.
H. Zhao , G. Moschetti , S. Wang , P-Å. Nilsson , and J. Grahn.
MoP06
MoP07
MoP8
Monolithic integration of InP based heterostructures on
silicon using SrTiO3 templates.
A. Chettaoui, B. Gobaut, J. Penuelas, J. Cheng1, G. Niu, L.
Largeau, P. Regreny, G. Saint-Girons.
Crystal structure X-ray investigation of InAs nanorods on
Si(111).
A. Davydok , A. Biermanns , M. Dimakis, S. Breuer, L. Geelhaar,
and U. Pietsch.
Structural and optical properties of InN films grown on
ZnO(000-1) by plasma-assisted molecular beam epitaxy.
Y. J. Cho, O. Brandt, M. Ramsteiner, M. Wienold and H. Riechert.
MoP09
Multifunctional Epitaxial Nanocomposite Films by LMBE.
J.Xiong.

POSTER SESSION 1
Monday, March 21 st 19:00 – 20:45
OPENING SESSION
MoP10
MoP11
MoP12
MoP13
MoP14
Near infrared high efficiency InAs/GaAsSb QDLEDs: band
alignment and carrier recombination mechanisms.
A. Hierro, M. Montes, M. Moral, J.M. Ulloa, A. Guzman.
In-situ Reflectance Anisotropy Spectroscopy (RAS) for
doping control during MBE growth of AlGaInAsSb laser
structures .
D. Hoffmann , T. Loeber and H. Fouckhardt.
Effect of growth temperature on surface morphology of
selectively grown GaN layers by ammonia-based metalorganic
molecular beam.
S. Naritsuka , C. H. Lin, R. Abe, S. Uchiyama, Y. Uete
and T. Maruyama .
Investigation of the local electronic structure of Cu-doped
GaN grown by plasma assisted MBE.
R. Schuber, P.R. Ganz , F. Wilhelm , A. Rogalev ,
and D.M. Schaadt.
Growth Optimization for InAs/GaSb T2SL Structures by
MBE.
Y. X. Song , S. M. Wang, C. Asplund, H. Malm, X. Lu, J. Shao.
MoP15
MoP16
A prototype of heterovalent interfaces: Reduction of the
potential barrier in the conduction band at the n-ZnSe /
n-GaAs in.
A. Frey, U. Bass, S. Mahapatra, C. Schumacher, J. Geurts
and K. Brunner.
Effect of growth temperature on quantum dot laser (Ga,In)
(N,As) self-assembled quantum dots.
O. A. Niasse, M. AL Khalfioui, B. Ba, A. Bèye, M. Leroux.
MoP17
MoP18
Effect of different monolayer coverage for the seed layer in
quaternary alloy capped multilayer InAs/GaAs quantum
dot system.
S. Chakrabarti ,A. Mandal and N. Halder .
Neutron reflectometry studies of hetero-interfacial H layer
in highly lattice-mismatched epitaxy on Si.
H. Asaoka, T. Yamazaki, D. Yamazaki, M. Takeda
and S. Shamoto.

POSTER SESSION 1
Monday, March 21 st 19:00 – 20:45
OPENING SESSION
MoP19
MoP20
Surface Electronic Properties of GaAs Nanowires.
O. Demichel, M. Heiss, J. Bleuse , H. Mariette, A. Fontcuberta.
Critical thickness of 2D-3D and “hut”-“dome” transitions
at the growth of GexSi1-x and Ge/GexSi1-x layers on the
Si(100).
V.A. Timofeev , A.I. Nikiforov , V.V. Ulyanov , O.P. Pchelyakov.
MoP21
Epitaxy and characterization of GaMnAs.
Martin Utz , D. Schuh , D. Bougeard and W. Wegscheider.
MoP22
MoP23
MoP24
MoP25
MoP26
MoP27
Mid-infrared Quantum Dot LEDs and microdisk laser grown
by MBE
A. Hochreiner, M. Eibelhuber, T. Schwarzl, H. Groiss, V.
Kolkovsky, G. Karczewski, T. Wojtowicz, W. Heiss, G. Springholz.
High-quality structures of InAs QDs in Al0.9Ga0.1As
matrix grown by droplet epitaxy.
A.A.Lyamkina, D.S. Abramkin, D.V.Dmitriev, S.P.Moshchenko,
. T.S. Shamirzaev, A. I. Toropov,K. S. Zhuravlev.
Structural properties of InAlN single layers nearly latticematched
to GaN grown by plasma assisted molecular beam
epitaxy.
Ž. Gacevic1, S. Fernández-Garrido and E. Calleja, D. Hosseini,
S. Estradé and F. Peiró.
Composition studies of site-controlled quantum dots.
G. Biasiol, V. Baranwal , S. Heun, M. Prasciolu, M. Tormen,
A. Locatelli, T. O. Mentes, M. N. Orti, and L. Sorba.
Single InAs quantum dots morphology and local electronic
properties on (113)B InP substrate.
C. Cornet , P. Turban , N. Bertru , S. Tricot , O. Dehaese
and A. Le Corre.
Investigations of growth kinetics of InN using pulsed RF
MBE.
A. Kraus , R. E. Buß, H. Bremers, U. Rossow, and A. Hangleiter.

POSTER SESSION 1
MoP28
MoP29
MoP30
MoP31
MoP32
MoP33
MoP34
MoP35
MoP36
Monday, March 21 st 19:00 – 20:45
OPENING SESSION
Low Thermal Budget Fabrication of Local Artificial
Substrates by Droplet Epitaxy on Silicon.
S.Bietti, C.Somaschini, N.Koguchi and S.Sanguinetti.
Investigation of improvement and degradation of thermal
annealed indium rich InGaN/GaN quantum wells grown by
NH3-MBE.
N. A. K. Kaufmann, A. Dussaigne , D. Martin and N. Grandjean .
Shape Changes in Patterned Planar InAs as a Function of
Thickness and Temperature.
K.G. Eyink , L. Grazulis , K. Mahalingham , M. Twyman , J. Shoaf,
V. Hart , J. Hoelscher , C. Claflin, and D. Tomich.
Influence of Al on the group III-assisted growth of axial
AlGaAs/GaAs heterostructure nanowires.
T. Rieger , M. I. Lepsa , H. Lüth , T. Schäpers
and D. Grützmacher.
Ferromagnetic and transport properties of very thin
(Ga,Mn)As layers.
L. Ebel, F. Greullet, T. Naydenova, J. Constantino, S. Mark,
C. Gould, K. Brunner and L.W. Molenkamp.
Self-assembled InP-nanoneedles grown on (001) InP by
gas source MBE.
M. Chashnikova, V. Bryksa, A. Mogilatenko, O. Fedosenko, S.
Machulik, M.P.Semtsiv, W. Neumann, and W.T. Masselink.
Selective growth of InP on pre-patterned wafers by the
means of Gas-Source MBE.
A.Aleksandrova, G.Monastyrskyi, O.Fedosenko, M.Chashnikova ,
S.Machulik ,J.Kishkat , M.P.Semtsiv and T.W.Masselink.
Scaling of quantum cascade laser efficiency with a number
of cascades.
O.Fedosenko, A.Aleksandrova , G.Monastyrskyi, M.Chashnikova ,
S.Machulik ,J.Kishkat, M.Klinkmüller, M.P.Semtsiv,T.W.Masselink.
MBE growth of quantum-cascade laser on pre-patterned
substrates.
G. Monastyrskyi, O. Fedosenko, M. Chashnikova, A. Alexandrova,
. S. Machulik,J.Kischkat, M. Klinkmüller, M. P. Semtsiv
and W. T. Masselink.

POSTER SESSION 1
MoP37
Monday, March 21 th -- 19:00 – 20:45
OPENING SESSION
Non polar GaN/ZnO heterostructures grown by ammonia
source molecular beam epitaxy.
J. Brault, G. Sophia, S. El Kazzi, J.-M. Chauveau, P. Vennéguès,
M. Nemoz, M. Teisseire, M. Leroux, C. Deparis, C. Morhain,
O. Tottereau, L. Nguyen.
MoP38
MoP39
Emission of colloidal nano-crystals embedded in MBE
grown ZnSe microstructures.
J. Kampmeier, M. Rashad , A. Pawlis, D. Schikora, K. Lischka.
Reproducible temperature calibration technique for GaAs
MBE.
François Morier-Genoud and Denis Martin.

POSTER SESSION 2
Tuesday, March 22 nd 19:00 – 20:45
TuP01
TuP02
OPENING SESSION
Fabrication and optical properties of CdTe quantum dots in
ZnTe nanowires.
P. Wojnar, E. Janik , A. Petroutchik , L. Baczewski , M. Goryca ,
T. Kazimierczuk , P. Kossacki , G. Karczewski and T. Wojtowicz.
In-situ, real time Auger Monitoring as a new tool for
growth characterization and control.
P. Staib.
TuP03
Optimisation of Unusual Quantum Dot Growth Conditions
for Optical Coherence Tomography Applications.
M. Hugues , M. A. Majid , S. Vezian , D. T. D. Childs , K. Kennedy
. and R. A.Hogg.
TuP04
A New Route for Strain relaxation in In0.52Al0.48As on
GaAs Grown by MBE.
S. M. Wang ,Y. X. Song , Z. H. Lai, M. Sadeghi and J. R. Dong.
TuP05
Nonpolar III-nitride microcavities for polariton lasing.
A. Dussaigne, G. Rossbach, J. Levrat, H. Teisseyre, I. Grzegory,
R. Butté, TSuski, and N. Grandjean.
Annealing effects on site-selective InAs quantum dots.
TuP06
TuP07
M. Helfrich , J. Hendrickson , M. Gehl , D. Rülke , D. Z. Hu, M.
Hetterich ,S. Linden , M. Wegener , H. Kalt , G. Khitrova ,
H. M. Gibbs and D. M. Schaadt.
Comparison of InAs quantum dots grown by ripening on
InP and GaInAsP buffer layers on InP(001).
P. Regreny, A. Benamrouche, C. Brillard and M. Gendry.
TuP08
TuP09
Study on homoepitaxial germanium nanowire growth.
J. Schmidtbauer, R. Bansen, T. Boeck, R. Heimburger,
Th. Teubner and T. Schoeder.
Correlating electronic and structural properties of Gaassisted
GaAs nanowires via cathodoluminescence
imaging.
J. Kasprzak,J.-S. Hwang, F. Donatini, C. Bougerol, H. Mariette,
Le Si Dang and R. Songmuang.

POSTER SESSION 2
Tuesday, March 22 nd 19:00 – 20:45
OPENING SESSION
Growth and microstructure of GaN:Cu.
TuP10
TuP11
TuP12
P. R. Ganz , G. Fischer , C. Sürgers, H. T. Hsing , L. Chang
And D. M. Schaadt.
Different strategies towards the deterministic coupling of
a Single QD to a Photonic Crystal Cavity Mode.
J.Herranz, I.Prieto, Y.González, J.Canet‐Ferrer, P.A.Postigo,
B.Alén, L.González, L.J.Martínez, M.Kaldirim, D. Fuster,
G.Muñoz‐Matutano, and J.Martínez‐Pastor.
Capping effect on the morphological and optical properties
of GaAs/AlGaAs quantum structures.
M. Jo, G. Duan, T. Mano and K. Sakoda.
TuP13
Optical signatures of dopant complexes in Arsenic doped
HgCdTe epilayers.
F. Gemain , I. C. Robin , B. Polge and A. Lusson.
TuP14
TuP15
Optical properties of post-growth annealed type-II GaSb
quantum dots.
A. Schramm, V. Polojärvi, T. V. Hakkarainen, A. Gubanov,
J. Paajaste, R.Koskinen, S. Suomalainen, and M. Guina.
Metamorphic 6.3Å GaInSb templates grown on GaAs
substrates for mid-infrared lasers.
L.Cerutti , J.B. Rodriguez and E. Tournié.
TuP16
Reflection high-energy electron diffraction phi scans forthe
in-situ monitoring of the growth of GaN nanowires on Si.
P. Dogan , O. Brandt, L. Geelhaar, and H. Riechert.
TuP17
TuP18
Optical polarization from self-organized InP QDs grown on
an self-undulated template.
A. Ugur, F. Hatami, N. Vamivakas L.Lombez ,M. Atatüre
B. and W. T. Masselink.
The MBE growth of HgCdTe on CdZnTe and CdTe/Ge at
CEA-LETI.
Giacomo Badano, Philippe Ballet, Sebastien Renet,
Philippe Duvaut, Xavier Baudry,Bernard Polge and Alain Million.

POSTER SESSION 2
Tuesday, March 22 nd 19:00 – 20:45
TuP19
OPENING SESSION
Control of nitrogen plasma for growth of GaN by plasmaassisted
MBE.
Z.R. Zytkiewicz, M. Sobanska, K. Klosek, H. Teisseyre, A.
Wierzbicka, E. Lusakowska, and W. Jung.
GaN/AlN semipolar quantum dots for ultra-violet emission.
TuP20
TuP21
A. Kahouli , N. Kriouche , J. Brault , B. Damilano , P. de Mierry ,
. A. Courville , J.Massies.
MBE growth approaches for improving Sb-based
In0.5Ga0.5As(Sb)/GaAs QDs.
M.J. Milla , Á. Guzmán, J.M. Ulloa, A. Hierro.
TuP22
Post-growth rapid thermal annealing of InAs quantum
dots grown on GaAs nanoholes formed by droplet epitaxy.
B. Alén, L. Wewior, D. Fuster, L. Ginés, Y. González, J. M.
Llorens, D. Alonso-Álvarez, and L. González.
TuP23
Ga blocking effect for GaN growth with NH3.
B. Damilano , A. Kahouli , J. Brault , D. Lefebvre , and J. Massies.
TuP24
Use of RHEED to optimize atomic layering of complex
Oxides.
B. A. Davidson, A. Yu. Petrov and S. Nannarone.
TuP25
TuP26
TuP27
II-VI nanostructures, with type-II band alignment, for
photovoltaics.
R. André, E. Bellet-Amalric, J. Bleuse, C. Bougerol,
M. Den Hertog, L. Gérard, H. Mariette.
Enhanced intermixing in Ge nano-prisms on groove
patterned Si (1 1 10) substrates.
G. Chen, G. Vastola, J. J. Zhang, B. Sanduijav, G. Springholz,
W. Jantsch, F. Schäffler.
Influence of nitrogen flux on InGaN growth by PAMBE.
H. Turski , M Siekacz , M. Sawicka , G. Cywiński , M. Kryśko ,
S. Grzanka ,I. Grzegory , S. Porowski , Z. Wasilewski
and C. Skierbiszewski.

POSTER SESSION 2
TuP28
TuP29
TuP30
Tuesday, March 22 nd 19:00 – 20:45
OPENING SESSION
Growth of metal Co/Ag superlattices on MgO(001):
microstructure and magnetic characterization.
Ana Ruiz, Enrique Navarro, María Alonso, Pilar Ferrer,
Daniel Margineda, F. Javier Palomares, Federico Cebollada,
Jesús Mª González.
Quantum dot formation from sub-critical InAs layers
grown on metamorphic InGaAs .
P. Frigeri , G. Trevisi and L. Seravalli.
Reversible Nanofacetting and 1D Ripple Formation of Ge
on High-Indexed Si (11 10) Substrat.
G. Springholz, B. Sanduijav and D. Matei.
TuP31
TuP32
TuP33
TuP34
TuP35
Homoepitaxy and nitrogen doping
of non polar ZnO films.
D. Tainoff, J.-M. Chauveau, C. Deparis, B. Vinter, M. AlKhalfioui,
M. Teisseire, Christian Morhain.
Structural characterization of MBE grown GaP/Si
nanolayers.
A. Létoublon, W. Guo, G. Elias, C. Cornet, A. Ponchet, A. Bondi,
T. Rohel, N.Bertru, C. Robert, T. Nguyen Thanh, O. Durand,
J.S. Micha and A. Le Corre.
An Auger Electron Analyzer System for In situ MBE
Stoichiometry Control.
W. Laws Calley, Phillipe Staib, Jonathan Lowder,
and W. Alan Doolittle.
MBE droplet epitaxy of InGaAs/Ga (100): effect of MBE
conditions and post-annealings on (Ga,In)droplet and
GaInAs nanostructure morphology and emission.
Chantal Fontaine, Poonyasiri Boonpeng, Guy Lacoste, Alexandre
Arnoult, Hejer Makhloufi, Olivier Gauthier-Lafaye, Guilhem
Almuneau, Somchai Ratanathammaphan, Somsak Panyakeow.
Synthesis of AlGaN nanowires by Molecular Beam Epitaxy.
A. Pierret , C. Bougerol , B. Gayral , B. Attal-Trétout ,
A. Loiseau and B.Daudin .
TuP36
Correlation of structural, chemical and optical characterization
of CdSe quantum dots inserted in ZnSe nanowires
M. Elouneg-Jamroz, M. den Hertog, S. Bounouar, E. Bellet-
Amalric, R. André,Y. Genuist, K. Kheng, J-P Poizat, S. Tatarenko.

POSTER SESSION 2
Tuesday, March 22 nd 19:00 – 20:45
OPENING SESSION
Optical and structural properties of InGaN/GaN nanowires.
TuP37
G. Tourbot , C. Bougerol, C. Leclere, B. Gayral, P. Gilet,
H. Renevier and B. Daudin.
TuP38
Position controlled self-catalyzed growth of GaAs
nanowires by molecular beam epitaxy.
Andreas Rudolph, Joachim Hubmann, Markus Kargl,
Benedikt Bauer, Marcello Soda, Josef Zweck, Dieter Schuh,
Dominique Bougeard and Elisabeth Reiger.

TERRITORY
Silicon Wafers, Ultrapure Silicon Dopant MBE
GaAs wafers FREIBERGER
EUROPE except Germany,Austria
InP wafers INPACT EUROPEAN research centers
GaSb,GaP,,InAs,InSb
SIMS and RBS Analysis Service PROBION
ARSENIC 7N,7N5 MBE slugs, FURUKAWA
GALLIUM 7N+, 7N5 +( RRR > 75000) MBE
ALUMINIUM 6N5 , the purest in the world
INDIUM 7N MBE
PHOSPHOROUS 7N MBE,6 N RASA
BERYLLIUM 5N+ MBE
EUROPE
WORLD except France,Japan
EUROPE
EUROPE
WORLD
EUROPE
EUROPE
WORLD
Bi, Cd, Pb, Sb, Sn, Te, Zn, Se, Th, Mg6N, Mn 5N8, S 6 or 6N +GaTe,GeS pieces
PBN, PG crucibles,pieces
WORLD
WORLD
Ga Recycling, GaAs, InP Wafers Reclaim
Bonding Wires and Ribbons MEM.SUMITOMO METAL MINING
Bonding Caps-Wedges, Pick up tools , Die Collets, SPT ROTH
Western/Eastern EUROPE
FRANCE
Resinoid Blades THERMOCARBON
www.azeliselectronics.com
Tel : 0033-1-44 73 10 70 Fax : 0033-1-44 73 10 53 – Azelis.Electronics@Azelis.fr

Monday Session Mo1
Arsenides I

Mo1.1
GaAs based nanostructures grown by droplet epitaxy
S.Sanguinetti * , C. Somaschini § , S. Bietti and N. Koguchi
L-NESS and Dip. di Scienza dei Materiali, Università di Milano Bicocca, Via Cozzi 53, 20125 Milano, Italy
What makes three dimensional semiconductor quantum nanostructures (QN) so attractive is the
possibility to tune their electronic properties by careful design of their size and composition. These
parameters set the confinement potential of electrons and holes, thus determining the electronic and
optical properties of the QN. An often overlooked parameter, which has a even more relevant effect on
the electronic properties of the QN, is shape. Gaining a strong control over the electronic properties of
semiconductor nanostructure via shape tuning is the key to access electronic fine design possibilities.
We present an innovative growth method, the Dropled Epitaxy (DE) [1,2], a variant of molecular
beam epitaxy, for the fabrication of semiconductor III-V QNs with highly designable shapes and
complex morphologies. In short, the DE growth procedure consists of first irradiating the substrate
with a group III molecular beam flux, leading to the formation of numerous, nanometer-sized, metallic
droplets on the surface which are subsequently crystallized into nanostructures by a group V molecular
beam. With DE is possible to combine multiple single QNs, namely quantum dots, quantum rings
and quantum disks, with tunable sizes and densities, into a single multi-functional QN thus allowing
an unprecedented control over the electronic properties of the QNs [2,3] (see Figure 1). In addition,
DE is intrinsically a low thermal budget growth of III-V materials, being fully performed at 200-
350°C. This makes DE perfectly suited for the realization of growth procedures compatible with backend
integration of III-V materials on Si [4,5].
Fig 1: Three examples of possible complex quantum strcutures fabricated by Droplet Epitaxy: (left panel) double
concentric double rings, (central panel) triple concentric quantum rings and (right panel) coupled ring disks.
[1] N. Koguchi, et al., J. Cryst. Growth (1991), 111, 688
[2] C. Somaschini, et al., Nano Letters 2009, 9, 3419
[3] C. Somaschini, et al., Nanotechnology 2010, 21, 125601
[4] S. Bietti, et al., Appl. Phys. Lett. 2009, 95, 241102
[5] C. Somaschini, et al. Appl. Phys. Lett. 2010, 97, 053101
* Corresponding author: stefano.sanguinetti@unimib.it
§ Present Address: Paul Drude Institut, Berlin (D)

Mo1.2
Interfacial strains in InAs/AlSb multilayers for short wavelength
quantum cascade lasers
C. Gatel 1,3 , B. Warot-Fonrose 1,3 , A. Ponchet 1,* , C.Magen 2,3 , R.Ibarra 2,3 , R. Teissier 4
and A.N. Baranov 4
1
CEMES-CNRS, 29 rue Jeanne Marvig 31055 Toulouse, France
2 INA, Universidad de Zaragoza, C/ Mariano Esquillor, Edificio I+D, 50018 Zaragoza, Spain
3
TALEM Associated Laboratory CNRS-University of Zaragoza, 29 rue Jeanne Marvig 31055 Toulouse, France
4 IES-UMR 5214, CNRS-Université Montpellier 2, 34095 Montpellier, France
This work explores the interfacial strain in InAs/AlSb multilayers grown by MBE on InAs(001)
substrates. Due to the remarkably high discontinuity of the conduction band (2.1 eV), InAs/AlSb
constitutes a highly interesting system for quantum cascade lasers (QCLs). These devices can cover
the wavelength region from 2.7 to 4.0 µm, well beyond the limits of QCLs based on the well matured
InP system [1,2].
In InAs/AlSb based QCLs, the wavelength of emission mainly depends on the layers thicknesses,
which have to be controlled at a fraction of monolayer. In addition, to allow the tunnel effect between
adjacent InAs wells, the AlSb barriers can be locally as thin as 4 or 5 atomic planes. So, the interfacial
zones cannot be neglected. The well and the barriers do not present common atoms: the interfaces
consist of either AlAs or InSb bonds. Consequently, while the average lattice mismatch is low (1.3%),
interfaces are susceptible to concentrate high lattice distortion. Indeed, in the bulk state AlAs and InSb
presents lattice mismatch with InAs of -7% and 7%, respectively. Schematically, the interfaces can be
tensile (Al-As type), compressive (In-Sb type) or neutral (alloyed type).
We have grown different samples with various procedures which aim at tuning the interfacial zones. In
a first sample, so-called reference sample, the procedure consisted in a simple growth interruption of 3
seconds under As flux for the AlSb on InAs interfaces, under Sb flux for the InAs on AlSb interfaces.
In other samples, we modified the growth sequences to intentionnaly favour either the Al-As or the In-
Sb type interfaces, by modifying the group V flux and/or the duration of growth interruptions. The
samples were examined by high resolution electron microscopy (HREM) on a TECNAI F-20 equipped
with a corrector of spherical aberration, in order to analyse the out-of-plane strain profiles [3,4].
Fig.1 (reference sample) shows a relatively symmetric profile at the two interfaces. The negative outof-plane
lattice distortion suggests a moderate tensile stress, probably due to the formation of Al-As
type interfaces. It has been possible to modify qualitatively and quantitatively this profile. An example
is given in fig. 2, where a strong negative strain is achieved at the AlSb on InAs interface; this
suggests the formation of a tensile interfacial zone of the Al-As type. A strong positive strain is
achieved at the reverse interface; this suggests the formation of a compressive interfacial zone mainly
of the In-Sb type. Chemical mapping using high angle annular dark field (HAADF) performed on a
probe-corrected STEM (fig.3) evidences an asymmetry of the two interfaces in agreement with the
hypothesis of an Al-As type interface for the AlSb on InAs interface.
To summarize, interfacial strains in InAs/AlSb multilayers have been studied. High compressive and
tensile stresses have been revealed at interfaces between InAs and AlSb in some samples. Specific
MBE growth sequences were used to control composition and strain of interfacial zones.
[1] J. Devenson, D. Barate, R. Teissier, and A. N. Baranov, Electronics Letters, 42, 1284-1286 (2006).
[2] J. Devenson, O. Cathabard, R. Teissier, and A. N. Baranov, Applied Physics Letters, 91, 141106 (2007)
[3] M.J. Hÿtch, E. Snoeck, and R. Kilaas, Ultramicroscopy 74, 131 (1998)
[4] C. Gatel et al, Acta Mat., 58, 3238 (2010)
[5] This research was supported by C’NANO GSO (CNRS, France)
* Contact: anne.ponchet@cemes.fr

Mo1.2
-3% +3%
(a)
(b)
Fig 1: Reference sample (InAs and AlSb thicknesses are 20 nm and 4 nm, respectively).
(a) HREM image in zone axis, (b) map of variation of the out-of-plane lattice parameter (the reference is
InAs), and profile of this variation across the small window.
-6% +6%
(a)
(b)
Fig 2: Same than fig.1 in a modified sample where Al-As and In-Sb interfaces have been favored by the growth
sequences adopted at interfaces.
(a)
(b)
Fig 3: STEM HAADF image of the same sample than in fig.2 (a) general view, (b) detail.
The AlSb on InAs interface is characterized by a darker contrast than InAs and AlSb. This is the signature of a light
material compared with InAs and AlSb, in agreement with the hypothesis of an Al-As type interface.

Mo1.3
200 mm GaAs wafers by MBE on SGOI and Ge/Si substrates
M. Richter 1,* , T. Topuria 2 , C. Marchiori 1 , M. El-Kazzi 1 , C. Rossel 1 , C. Gerl 1 , D.J. Webb 1 ,
T. Smets 3 , C. Andersson 1 , M. Sousa 1 , D. Caimi 1 , L. Czornomaz 1 , H. Siegwart 1 ,
J.-F. Damlencourt 4 , J.-M. Hartmann 4 , P.M. Rice 2 , and J. Fompeyrine 1
1
IBM Research – Zurich, Säumerstrasse 4, 8803 Rüschlikon, Switzerland
2 IBM Research – Almaden, 650 Harry Road, San Jose, California, USA
3 Kath. Uni. Leuven – Celestijnenlaan 200D, 3001 Heverlee, Belgium
4 CEA-LETI – Grenoble, 17 rue des Martyrs, 38054 Grenoble, France
The use of production size wafers is a prerequisite for the integration of III-V materials in future
CMOS technology. Besides obvious technical benefits to merge III-V with silicon process technology,
it also allows to economize on scarce materials. In this presentation, GaAs heterogeneous integration
on 200 mm Si and SGOI wafers via Ge buffers will be discussed.
In our previous work, we have studied thin (≤ 250 nm) MBE-grown Ge buffers. This had the
advantage to rely on the use of two UHV-connected MBE chambers, thus limiting air exposure
between Ge buffer and GaAs growth [1]. In this paper, we discuss an alternative approach which
consists in combining Ge growth by reduced pressure-chemical vapor deposition (RP-CVD) with
subsequent, ex-situ MBE GaAs deposition. CVD growth allows for thicker Ge layers and anneals in
the H 2 carrier gas. By this means reduced defect densities [2] and flatter surfaces due to the surfactant
action of H 2 [3] can be achieved. In a first series of samples, we use 200 mm Si(001) wafers with 1.5
µm Ge. In a second series of sample, we use 200 mm SGOI wafers, which feature improved insulation
to the Si substrate and optional co-integration of fully depleted GOI p-FETs [4].
Special care needs to be taken to clean the Ge/Si or SGOI wafers before the GaAs MBE growth. The
anneal temperature should be minimized to take account of the limited temperature stability. All
wafers were prepared with an HF-dip as last cleaning step. Then, they were either exposed to a remote
hydrogen plasma [5] or annealed at 600 °C in UHV. The removal of oxides and carbon was monitored
with x-ray photoelectron spectroscopy (XPS). The H plasma efficiently removes oxide species at
temperatures as low as 250 °C on both Ge/Si and SGOI substrates (spectra in Fig. 1). Only traces of C
contamination are observed after such cleaning treatment. As opposed to that, after sample flash
substantial amounts of SiOx are observed (Fig. 1(b)).
For the SGOI wafers, either 200 nm Ge were deposited in a UHV-connected Si/Ge MBE system or
SGOI wafers capped with 20 nm CVD-grown Ge were used. Then, the wafers were directly
transferred under UHV conditions for epitaxy of 500 nm GaAs. Fig. 2 shows corresponding reflection
high energetic electron diffraction (RHEED) images after the individual steps.
We will discuss the impact of the substrate, its cleaning and of the different buffers on the GaAs
structural quality. Cross-sectional transmission electron microscopy (XTEM) and atomic force
microscope (AFM) images of a first sample of 500 nm GaAs on 20 nm Ge CVD-capped SGOI are
shown in Fig. 3 (a) and (b), respectively.
TEM specimen preparation by L.M. Clark and L.E. Krupp (IBM Research - Almaden) as well as
financial support by the European Commission in frame of the FP7 project DUALLOGIC is gratefully
acknowledged.
[1] M. Richter, C. Rossel, D.J. Webb, T. Topuria, C. Gerl, M. Sousa, C. Marchiori, D. Caimi, H. Siegwart, P.M. Rice and J.
Fompeyrine, “GaAs on 200 mm Si wafers via thin temperature graded Ge buffers by molecular beam epitaxy” submitted to J.
Cryst. Growth and abstract MBE 2010 conference.
[2] J.M. Hartmann, A. Abbadie, N. Cherkashin, H. Grampeix, and L. Clavelier, Semicond. Sci. Technol., 24, 055002 (2009).
[3] L. Colace, G. Masini, F. Galluzzi, G. Assanto, G. Capellini, L. Di Gaspare, E. Palange, and F. Evangelisti, Appl. Phys.
Lett., 72, 3175 (1998).
[4] W. Van Den Daele, E. Augendre, K. Romanjek, C. Le Royer, L. Clavelier, J.-F. Damlencourt, E. Guiot, B. Ghyselen, and
S. Cristoloveanu, ECS Trans., 19, 145 (2009).
[5] C. Marchiori, D. J. Webb, C. Rossel, M. Richter, M. Sousa, C. Gerl, R. Germann, C. Andersson, and J. Fompeyrine, J.
Appl. Phys., 106, 114112 (2009).
__________________________
* Contact: mri@zurich.ibm.com

Mo1.3
(a)
(b)
Fig 1: XPS after the different cleaning steps of (a) Ge/Si and (b)SGOI wafers.
(a)
(c)
(b)
Fig 2: RHEED image of 200 mm Ge/Si wafer, (a) as loaded, (b) after Hydrogen clean, (c) after 500 nm GaAs growth.
(a)
(b)
Fig 3: 500 nm GaAs on 20 nm CVD-grown Ge on SGOI, (a) XTEM image, (b) AFM image.

Mo1.4
Tuning the size, strain and band offsets of InAs/GaAs quantum
dots through a thin GaAs(Sb)(N) capping layer
J.M. Ulloa 1,* , M. Montes 1 , K. Yamamoto 1 , A. Guzman 1 , A. Hierro 1 , M. Bozkurt 2 ,
P.M. Koenraad 2 , D. Fernández 3 , D. González 3 and D. Sales 3
1 ISOM and Dpto. Ing. Electronica, Univ. Politecnica Madrid, Ciudad Universitaria s/n, 28040 Spain
2 COBRA-PSN, Eindhoven University of Technology, P.O. Box 513, NL-5600MB Eindhoven, The Netherlands
3 Departamento de Ciencia de los Materiales e I. M. y Q. I., Facultad de Ciencias, Universidad de Cádiz, Spain
A common approach used in the past few years to extend the emission wavelength of InAs/GaAs
quantum dots (QD) to the 1.3-1.55 µm range is the use of a strain reducing layer. By using thin
GaAsSb capping layers, photoluminescence (PL) emission at 1.55 µm has been demonstrated [1-3].
Nevertheless, adding Sb to the capping layer does not only affect the QD strain, but also the
QD/capping layer valence band offset, giving rise to a type-II band alignment at high Sb contents (~
17 %). The result is that emission at long wavelengths can only be achieved in these structures with a
type-II band alignment that degrades the PL. On the other hand, introducing small amounts of N in the
capping layer would strongly reduce the QD/capping layer conduction band offset, inducing an extra
red shift. This could allow reaching 1.55 µm while keeping a type-I band alignment. In this work, we
show how the height, strain and QD/capping layer band offsets can be tuned by using a thin
GaAs(Sb)(N) capping layer, allowing to reach emission in the 1.55 µm region. A detailed analysis
by means of PL, cross-sectional scanning tunneling microscopy, transmission electron
microscopy and atomic force microscopy allows correlating the optical and structural
properties of the QD structures.
The analyzed samples were grown by solid source MBE on n+ Si doped (100) GaAs
substrates. In all of the samples 2.7 monolayers (ML) of InAs were deposited at 450 °C and
0.035 ML/s on an intrinsic GaAs buffer layer. The QDs were capped with a nominally 5 nm
thick GaAs(Sb)(N) layer grown at 470 °C. The nominal Sb and N contents were changed
from 0 to 30 % and 0 to 3 %, respectively.
Besides the expected reduction in the QD strain and the transition to a type-II band alignment
at high Sb contents, we find that the QD height can be controlled through the amount of Sb in
the capping layer due to reduced In-Ga intermixing during the capping process [4]. Figure 1
shows how the QD height increases progressively with the amount of Sb, reaching the value
of uncapped QDs for ~ 13% Sb. The increased height leads to improved PL at moderate Sb
contents compared to the reference GaAs-capping sample (Fig. 2). On the other hand, GaAsN
capping layers strongly reduce the QD conduction band offset inducing also a PL red shift but
preserving the type-I alignment and the optical quality (see Fig. 2). The impact of N on the
structural properties of the QDs will be discussed. The longest wavelengths are achieved with
the simultaneous presence of both Sb and N in the capping layer (Fig. 3), which allows to
independently modify the conduction and valence band offsets. Nevertheless, in this case the
PL spectra are degraded compared to the case of the counterpart ternary alloys (Fig. 2).
Although a type-I band alignment is preserved in this case even at the longest wavelengths,
increasing the N content in the GaAsSb capping layer induces a progressive degradation of
the PL spectrum (Fig. 3). The reasons for this are investigated and will be discussed.
[1] K. Akahane et al., Physica E (Amsterdam) 21, 295 (2004).
[2] J. M. Ripalda et al., Appl. Phys. Lett. 87, 202108 (2005).
[3] H. Y. Liu et al., J. Appl. Phys. 99, 046104 (2006).
[4] J.M. Ulloa et al., Phys.Rev.B. 81,165305 (2010).
_____________
* Contact: jmulloa@die.upm.es

Mo1.4
Capped QD height/Uncapped QD height
1.1
1.0
0.9
0.8
0.7
0.6
0.5
0.4
0.3
GaAsSb
Reduced
decomposition
Completely
supressed
decomposition
0 3 6 9 12 15 18 21 24
Sb content (%)
Fig 1: QD height normalized to the height of the uncapped
QDs as a function of the Sb content. A value of 1.0
indicates a completely suppressed decomposition process.
The two different background colors indicate two different
regimes regarding QD decomposition.
250
200
GaAsSb
(~13%)
15 K
PL intensity (a.u.)
150
100
50
GaAs
GaAsN
GaAsSbN
x 10
0
1000 1100 1200 1300 1400
Wavelength (nm)
Fig 2: 15 K PL spectra of InAs/GaAs QDs grown under
the same conditions and capped with different alloys.
1,2
GaAsSbN
> N
15 K
Normalized PL intensity
1,0
0,8
0,6
0,4
0,2
0 % N
0,0
1000 1100 1200 1300 1400 1500 1600
Wavelength (nm)
Fig 3: Normalized 15 K PL spectra of a series of samples
with constant Sb and increasing N content in the capping
layer.

Mo1.5
InAs Quantum Dot Chains Grown on Nanoimprint Lithography
Patterned GaAs(100)
T. V. Hakkarainen*, A. Schramm, J. Tommila, A. Tukiainen, R. Ahorinta, M.
Dumitrescu, M. Guina.
Optoelectronics Research Centre, Tampere University of Technology,
P.O. Box 692, FIN-33101 Tampere, Finland.
The ability to fabricate Stranski-Krastanov quantum dots on pre-determined locations, i.e. sitecontrolled
growth, is essential for enabling emerging nanophotonic applications, such as photonic
integrated circuits incorporating quantum dot chains (QDCs) as nanophotonic waveguides. As a
method for the fabrication of site-controlled InAs QDCs, we combine growth by molecular beam
epitaxy and nanoimprint lithography (NIL) [1]. NIL is able to produce sub 10 nm linewidths with high
throughput and enables fast processing of large wafer area.
In this paper we focus on studying structural and optical properties of QDCs with varying orientations
with respect to the substrate crystal directions. The investigated samples were prepared in three stages.
In the first stage, a 100 nm GaAs buffer, a 100 nm AlGaAs layer, and a 100 nm GaAs were deposited
at 590 °C on n-GaAs(100) substrates by MBE. Then grooves were ex situ patterned by UV-NIL. The
groove width was 90 nm, depth 30 nm, and period 180 nm. In the final stage, the patterned surface
was covered with a 60 nm GaAs regrowth buffer at 490 °C and 2.2 ML InAs QDs grown at 515 °C.
For optical investigation, the QDs were covered with GaAs and AlGaAs layers. We show that this
method enables the simultaneous growth of QDCs oriented along [011], [01-1], [010], and [001]
directions (Fig. 1) exhibiting strong photoluminescence (PL) emission at room temperature. Being
able to form QDCs with different orientations with at the same growth conditions is crucial for the
fabrication of QDC networks for integrated circuits. Furthermore, we report low temperature PL
(temperature, power, and polarization dependencies) for the optical characterization of the QDCs and
atomic force microscopy (AFM) based facet analysis for investigating the morphology of the patterned
surface.
(a)
(c)
[011]
(b)
(d)
[01-1]
PL intensity (arb. units)
1.0
0.8
0.6
0.4
0.2
0.0
[0-11]
[011]
[0-1-1]
[01-1]
4.5 mW
1.2 mW
0.7 mW
0.1 mW
1.1 1.2 1.3 1.4
Energy (nm)
Fig. 1. AFM pictures of QDCs grown on a groove
pattern oriented along [011] (a), [01-1] (b), [010] (c),
and [001] (d) directions.The color scale in (a)-(d) is 26
nm.
Fig. 2. PL spectra from [01-1]-oriented QDCs measured at
10K with different excitation laser powers. The inset
shows polarization anisotropy of the PL emission. The
radial axis of the polar plot represents relative PL intensity
ranging from 0.9 to 1.1
[1] J. Tommila, A. Tukiainen, J. Viheriälä, A. Schramm, T. Hakkarainen, A. Aho, P. Stenberg, M. Dumitrescu and M. Guina,
”Nanoimprint lithography patterned GaAs templates for site-controlled InAs quantum dots”, to be published in J. Cryst.
Growth.
_________________________
* Contact: teemu.hakkarainen@tut.fi

Monday Session Mo2
Antimonides & Phosphides

Mo2.1
MBE growth of LiMnAs
Vít Novák
Institute of Physics of the Academy of Sciences, Cukrovarnická 10, 162 53 Praha, Czech Republic
Compound semiconductors derived from silicon have had a tremendous impact on the physics and
applications of semiconductors. Two textbook examples are the direct gap III-V semiconductors and
the prototype magnetic II-VI semiconductors, Fig.1. Remarkably, none of the other closest relatives of
silicon from the I-III-IV and I-II-V compounds have so far been synthesized by modern epitaxial
growth techniques and the potential of these compounds has remained virtually unexplored. In this
talk we focus on I-Mn-V compounds which surprisingly have not previously been considered as
candidate semiconductors. One of the key motivations to establish their semiconducting electronic
structure is that they are among the rare known silicon relatives with magnetic ordering temperature
safely above room temperature.
Fig.1: Closest relatives of silicon emerging by applying the “proton transfer” rule.
We demonstrate on LiMnAs that high-quality materials with group-I alkali metals in the crystal structure
can be grown by standard solid source molecular beam epitaxy [1]. The epitaxial LiMnAs film exhibits
optical gap, evidenced by optical transmission measurements and consistent with the band structure
obtained by our ab initio calculations. Squid magnetometry measurements support earlier reports of
high antiferromagnetic ordering temperature.
We propose a strategy for employing I-Mn-V compounds in high-temperature semiconductor
magneto-electronics. The key principle is to utilize relativistic magnetic and magneto-transport anisotropy
effects whose common characteristics is that they are an even function of the microscopic magnetic
moment vector and are therefore in principle equally well present in materials with ferromagnetic (FM) and
antiferromagnetic (AFM) order. The application of AFM semiconductors to exchange bias FMs opens an
immediate research opportunity for integrating conventional semiconductor micro and opto-electronics
functionalities directly in the exchange-biasing AFM layers in common magneto-electronic devices. In
these structures LiMnAs can be combined, e.g., with lattice matched FM semiconductor (In,Mn)As or
conventional transition metal FMs. The FM-AFM coupling can also be used for controlling the staggered
moment orientation in the AFM by the exchange spring effect induced by rotating moments in the
ferromagnet.
[1] T. Jungwirth, et al., Phys. Rev. B 83, 035321 (2011).
_______________________
* Contact: vit.novak@fzu.cz

Mo2.2
Growth and structural properties of InSb-based heterostructures
A.N. Semenov, B.Ya. Meltser, V.A. Solov'ev, T.A. Komissarova, Ya.V. Teren't'ev,
A.A. Sitnikova, and S.V. Ivanov
Ioffe Physical-Technical Institute of RAS, Polytekhnicheskaya 26, St. Petersburg 194021, Russia
Among III-V semiconductors, indium antimonide has the narrowest band gap and the smallest
effective electron mass that results in a record value of room-temperature electron mobility in bulk
InSb layer as well as 2D electron gas (2DEG). This makes InSb a very attractive material for new
generation of high speed devices like a high electron mobility transistor (HEMT) [1] and IR
photodetectors [2]. Al x In 1-x Sb alloys can be used as barriers for the InSb QW heterostructures but Al
content is limited by x~0.2 due to abrupt increase of the stress in high Al-content alloys grown on
InSb. This paper reports on MBE growth, structural, optical, and transport properties of InSb,
Al x In 1-x Sb (0≤x≤0.3) and their QW heterostructures.
All the structures were grown on semi-insulated GaAs substrates using RIBER 32P setup
equipped with conventional solid source effusion cells for Al, In and Sb 4 . Substrate temperature T S
was varied in the 380–550°С range and measured by an IR pyrometer calibrated using well-known
surface reconstruction transitions on GaAs, monitored in situ by reflection high energy electron
diffraction (RHEED). The mobility and carrier concentration were determined using Hall effect
measurements at temperatures 77 and 300 K. X-ray diffraction, transmission and scanning electron
microscopy (TEM and SEM), scanning probe microscopy and photoluminescence (PL) were
employed for structural and optical characterization of the InSb-based samples. PL spectra were
measured only for AlInSb alloys because the nitrogen-cooled InSb photodiode was used as a detector.
Due to the absence of InSb semi-insulated (SI) substrates the InSb/AlInSb 2DEG QW
heterostructures should be grown on strongly lattice-mismatched SI substrates of GaAs (Δa/a=14.6%).
To prevent the Sb evaporation induced by the giant mismatch we employed intermediate AlSb buffer
layer which is characterized by much higher dissociation temperature and lower Δa/a=7%. Moreover,
it was found that growth of Sb-based compounds on GaAs substrates is improved in case of exposure
of the GaAs surface to the Sb 4 flux. The RHEED pattern during this procedure changes from (2×4) for
As-stabilised surface to (2×8) one (Fig. 1) inherent to GaAs surface covered by a 1-ML-thick GaSb
layer [3]. This GaSb layer promotes transition to the 2D growth mode during initial stage of AlSb
growth. The perfect streaky (1×3)Sb RHEED pattern is observed after growth of 40 nm AlSb. The
initiation of the AlSb growth at low temperatures (480°С) leads to the formation of high density of
stacking faults and microtwins (Fig.2a). Such defects could lead to the strong anisotropy of transport
properties of 2DEG heterostructures [8]. The higher AlSb growth temperature (T S ~520°С) results in
suppression of microtwins formation at the AlSb/GaAs heterointerface while preserves the 1ML-GaSb
coverage. The microtwins formation then takes place only at the AlSb/AlInSb heterointerface grown at
low temperature inevitably (Fig. 2b), but their density is at least 5 times lower as compared with
heterostructures grown atop of the low-temperature AlSb layer.
Further optimization of the buffer structures, including inserting the thin InSb layers as well as
the InSb/AlInSb strained superlattices (SL), allowed us to reduce the dislocation density around the
top InSb QWs by two orders of magnitude. The room-temperature electron mobility and concentration
in 1-μm-thick InSb/GaAs heteroepitaxial layers were measured as 52 000 cm2/V×s and 3×10 16 cm -3 ,
respectively. In the case of AlInSb epilayers, concentration and mobility values were found to depend
strongly on Al content. Based on these findings, the InSb/AlInSb 2DEG QW heterostructures were
grown. The room-temperature electron concentration and mobility in the InSb channel were as high as
3.6×10 12 cm -2 and 20 500 cm 2 /V×s, respectively. We believe that optimization of the InSb/AlInSb
heterostructure design and improving the quality of heterointerfaces will enable one to improve
characteristics of InSb 2DEG.
This work is supported by the Russian Foundation for Basic Research Project #09-02-01500
[1] T. Ashley, L. Buckle, S. Dutta et al. Electron. Lett. 43, (2007) 777.
[2] Patrick J. Treado, Ira W. Levin, and E. Neil Lewis Applied Spectroscopy 48, (1994) 607.
[3] L.J. Whitman, B.R. Bennett, E.M. Kneedler et al. Surface Science 436 (1999) L707.
[4] J.B. Boos, B.R. Bennett, W. Kruppa et al. Sol. St. Electron. 47, 181 (2003)

Mo2.2
011
0ī1
Fig. 1. (2×8) RHEED pattern during antimony exposure of GaAs surface (T S = 480°C)
011
a
AlInSb
AlInSb
AlSb
GaAs
b
Fig. 2 Cross-section TEM images of GaAs/AlSb/AlInSb heterostructures: a) AlSb grown at
Ts ~ 480°Ñ, b) AlSb grown at Ts ~ 510°Ñ.
__________________________
* Contact: semenov@beam.ioffe.ru
AlSb

Mo2.3
Comparison of interface properties in two-dimensional
heterostructures grown layer-by-layer or by step flow
E. Luna 1* , R. Hey 1 , A. Guzmán 2 and A. Trampert 1
1
Paul-Drude Institut für Festkörperelektronik, Hausvogteiplatz 5-7, 10117 Berlin, Germany
2 ISOM-ETSI Telecomunicación-UPM, Ciudad Universitaria s/n, 28040 Madrid, Spain
By analyzing the experimental composition profiles obtained from transmission electron microscopy
(TEM) techniques, we have previously determined that the smooth variation of the element
concentration with position across the interface in different III-V heterostructures grown by molecularbeam-epitaxy
(MBE) follows a sigmoidal law: x(z)= x 0 /[1+exp(-z/L)], where the interface width L is
the main fitting parameter and x 0 as well as z denote the nominal mole fraction and the position along
the growth direction, respectively [1,2]. The wide range of material systems (group-III arsenides,
antimonides and III-V metastable quaternary compounds) exhibiting the sigmoidal dependence [2,3]
supports the hypothesis that it is a universal behavior, which is determined by fundamental processes
occurring during growth. In this respect, a key issue is to determine the role of the kinetics and to give
an answer to the always challenging question of the interplay between kinetics and thermodynamics.
In order to investigate the influence of the kinetics on the interface development, we consider
(In,Ga)As/GaAs heterostructures grown by MBE in the thermodynamically controlled Frank van der
Merwe (FM) two-dimensional (2D) growth mode under very different kinetic conditions. General
homoepitaxial growth experiments prove that, within the FM 2D mode, kinetic control is possible
through changes in the growth temperature, i.e., the adatom rate of diffusion, and/or the flux, i.e., the
rate of deposition. In particular, two distinct growth modes can be observed: step flow and layer-bylayer
growth by 2D island nucleation. We focus on 2D (In,Ga)As layers grown either by step flow or
by island nucleation. Thus, the samples would share the same thermodynamical constraints, but differ
largely in their kinetic aspects. The In distribution profile shown in Fig.1 corresponds to a 7 nm
In 0.4 Ga 0.6 As/GaAs quantum well (QW) grown by island nucleation on on-axis GaAs(001). As
observed, the change in composition across the interfaces is very well described by the sigmoidal law.
However, MBE growth on misoriented GaAs(111)B substrates provides a very good case study of
layers grown in the step flow mode. Fig. 2(a) shows a chemically sensitive g 002 dark-field (DF) TEM
image of a 7 nm In 0.3 Ga 0.7 As/GaAs QW grown on vicinal GaAs(111)B, where the growth proceeds via
step flow. The analysis of the experimental In composition profile shown in Fig.2(b) reveals that, in
spite of the broad and asymmetric interfaces, the change in composition across the interfaces also
follows a sigmoidal law. Thus, the extreme change in kinetics (step flow vs. island nucleation) does
not seem to modify the profile of the interface, which retains its sigmoidal dependence. These
experimental findings suggest that the general shape of the transition region at the interface is
determined by thermodynamics, although kinetic aspects are reflected in the interfacial width.
Further insight into the interplay between thermodynamics and kinetics can be gained from QWs
grown by the so-called solid phase epitaxy (SPE) [4]. In this approach, the (In,Ga)As QW is formed
after thermally induced crystallization of the respective amorphously deposited compounds [5]. Fig.3
represents the experimentally determined In profile from a SPE-grown (In,Ga)As/GaAs QW on
GaAs(001). As observed, both interfaces are remarkably well described by the sigmoidal function.
While exact processes occurring during SPE are not yet understood, obviously thermodynamics plays
a major role in the final step of the SPE method. This result further demonstrates the importance of
thermodynamic factors in the development of interfaces in III-V heterostructures.
[1] E. Luna, F. Ishikawa, P.D. Batista and A. Trampert, Appl. Phys. Lett. 92, 141913 (2008)
[2] E. Luna, F. Ishikawa, B. Satpati, J.B. Rodriguez, E. Tournié and A. Trampert, J. Cryst. Growth 311, 1739 (2009)
[3] E. Luna, B. Satpati, J.B. Rodriguez, A.N. Baranov, E. Tournié and A. Trampert, Appl. Phys. Lett. 96, 021904 (2010)
[4] N.G. Rudawski, K.S. Jones, S. Morarka, M.E. Law and R.G. Elliman, J. Appl. Phys., 105, 081101 (2009)
[5] R. Hey, P.V. Santos, E. Luna, T. Flissikowski and U. Jahn, J. Cryst. Growth (2010), doi:10.1016/j.jcrysgro.2010.09.020
__________________________
* Contact: luna@pdi-berlin.de

Mo2.3
In content (%)
40
30
20
10
0
Exp.
Sim.
growth direction
L lower
= 1.35 ML
L upper
= 1.35 ML
20 40 60 80
Position (ML)
Fig 1: Experimentally determined In profile for a (In,Ga)As/GaAs QW grown on GaAs(001) in the 2D
island nucleation mode. Solid line: simulated profile where the interfaces are defined by a sigmoidal
function. L lower and L upper denote the width of the lower and upper interface, respectively.
In content (%)
30
25
20
15
10
5
(b)
Exp.
Sim.
growth direction
L lower
= 2.6 ML
L upper
= 4.2 ML
0
-60 -40 -20 0 20 40 60
Position (ML)
Fig 2: (a) g 002 DF TEM micrograph of a (In,Ga)As/GaAs QW on GaAs(111)B, which has been grown in the step flow
mode. (b) Experimental In distribution and simulated profile, where the interfaces are defined by a sigmoidal function
30
Exp.
Sim.
L lower
= 1.8 ML
L upper
= 1.8 ML
In content (%)
20
10
growth direction
0
-20 0 20
Position (ML)
Fig 3: Experimentally determined In profile for a (In,Ga)As/GaAs QW grown on GaAs(001) by SPE
and simulated profile, where the interfaces are defined by a sigmoidal function.

Mo2.4
MBE growth of (GaAsPN/GaPN)/GaP quantum wells light
emitting diode
A. Bondi, C. Cornet * , W. Guo, O. Dehaese, N. Chevalier, C. Robert, T. Nguyen Thanh, S.
Richard, M. Perrin, A. Létoublon, J. P. Burin. J. Even, O. Durand and A. Le Corre
Université Européenne de Bretagne, France
INSA, FOTON, UMR 6082, F-35708 RENNES
Many recent research developments focus on the ability to integrate optical functions on
Silicon.[1] Among them, heterogeneous and coherent growth of lattice-matched GaPN alloy with ~2%
N on Si substrate should ensure defect-free heterostructures leading to long-life and stable optical
device on Si.[2] However, in these diluted nitrides devices, optical properties and band gap transitions
are always affected by the so-called clustering effects,[3] as well as random alloying effects.[4] While
the GaAsN material system has been widely studied,[5] GaPN and GaAsPN materials on GaP
substrate still need to be understood for the device improvement.[6-9]
Here, we present the growth of (GaAsPN/GaPN)/GaP(001) multi-Quantum Wells (QWs) light
emitting diode (LED) grown by gas source MBE. The LED structure is presented on Fig. 1. Growth is
performed at 580°C and 0.4µm/h in the whole structure. A 300 nm n-doped (1.10 18 cm -3 ) GaP:Si layer
is grown on GaP:S substrate. A 5 QW structure is then grown in the GaPN materials system, without
growth interruption. Nitrogen RF-plasma is set to ensure a x=0.005 composition into GaP 1-x N x . A p-
doped (2.10 18 cm -3 ) 300nm GaP:Be capping layer is finally grown. Details about thicknesses are
presented on figure 1(a). Electroluminescence (EL) is obtained at room temperature at 1.69 eV (734
nm) (fig. 1(b)). This EL spectrum is broad and likely composed of different contributions. In order to
understand this behavior, structural and optical studies are performed on both GaAsP/GaP and
GaAsPN/GaPN QWs. The 2 samples contain 5 QWs grown in the same conditions as for the LED.
An X-ray reflectometry (XRR) analysis is performed on both structures. The autocorrelation
function from the corrected XRR profile is presented on Fig. 2(a) and displays peaks centered on the
distances between interfaces, leading to the thicknesses and sum of contiguous thicknesses inside the
stacking. Similar shapes are obtained between the (GaAsPN/GaPN) and the (GaAsP/GaP) bilayer
period related peaks. A QW thickness of 2.7 ±0.1 nm is measured. A complementary X-ray diffraction
analysis is performed for both samples and allows for GaAsP/GaP superlattice an evaluation of the As
ratio with y=0.69±0.02 for the GaAs y P 1-y in the QWs. Then, assuming the same N incorporation in
both GaPN and GaAsPN, the N-based QW alloy composition is assumed to be GaAs 0.687 P 0.308 N 0.005 . In
the XRR, unexpected thickness contributions also appear. They are attributed to nitrogen or Asinduced
interfacial effects in the strain-compensated GaAsPN/GaPN/GaP system.
Then, a complete PL analysis is performed between 4K and 300 K. Evolution of the PL main
peak is shown on Fig. 2(b). PL emission is found to deviate from the Varshni law for semiconductors
when the temperature exceeds 150 K.[7,9] It also deviates from the S-shaped temperature dependence
observed especially in the GaAsN system.[5] A discussion is given on the origin of such temperature
dependence in GaAsPN in term of random alloy effects and clustering.
Finally, from these studies, some conclusions will be drawn on the future development of
GaAsPN/GaPN QWs lasers.
We acknowledge the Région Bretagne (CPER PONANT) for its financial support.
[1] D. Liang and J. E. Bowers, Nature Photon., 4, 511 (2010).
[2] H. Yonezu, Y. Kurukawa and A. Wakahara, J. Crystal Growth, 310, 4757 (2008).
[3] D. G. Thomas, J. J. Hopfield and C. J. Frosch, Phys. Rev. Lett., 15, 857 (1965).
[4] V. Popescu and A. Zunger, Phys. Rev. Lett., 104, 236403 (2010).
[5] R.J. Potter, N. Balkan, H. Cre, A. Arnoult, E. Bedel, and X. Marie, Appl. Phys. Lett., 82, 3400 (2003).
[6] I. A. Buyanova, G. Yu. Rudko, W. M. Chen, H. P. Xin and C. W. Tu, Appl. Phys. Lett., 80, 1740 (2002).
[7] A. Yu. Egorov, N. V. Kryzhanovskayaa, E. V. Pirogov and M. M. Pavlov, Semiconductors 44, 886 (2010).
[8] B. Kunert, K. Volz, J. Koch and W. Stolz, Appl. Phys. Lett., 88, 182108 (2006).
[9] C. Karcher, H. Grüning, M. Güngerich et al., Phys. Stat, Sol. (c), 6, 2638 (2009).
__________________________
* Contact: charles.cornet@insa-rennes.fr

Mo2.4
Fig 1: (a) structure of GaAsPN/GaPN light emitting diode, (b) Room temperature electroluminescence and
photoluminescence spectra measured from the GaAsPN/GaPN diode. Figure insert shows the I(V) characteristic.
Fig 2: (a) auto-correlation function from the X-ray reflectometry analysis for 5-stacked QWs including GaAsP/GaP
QWs and GaAsPN/GaPN QWs. Sample including Nitrogen shows additional interfaces. (b) Temperature dependence
of GaAsPN photoluminescence peak. Evolution is different from Varshni law and S-shaped dependence.

Mo2.5
Lattice mismatch accommodation at the GaSb/GaAs and
GaSb/GaP interfaces
L. Desplanque 1,* , S. El Kazzi 1 , C. Coinon 1 , Y. Wang 2 , P. Ruterana 2 , and X. Wallart 1
1
IEMN, CNRS and University of Lille UMR 8520, Avenue Poincaré – BP 60069 Villeneuve-d’Ascq
Cedex,FRANCE
2 CIMAP, CNRS-ENSICAEN-CEA-UCBN,6, Boulevard du Maréchal Juin, 14050 Caen Cedex, FRANCE
Antimony based compound semiconductors are gaining more and more interest since their unique
material properties (band offsets, electron and hole mobility) offer original solutions for electronic and
optoelectronic applications. The lack of large size and low cost lattice matched substrates for these
materials as well as the integration of these solutions in standard technological process require the
development of thin and high quality buffer layers accommodating the lattice mismatch with GaAs or
Si substrates. For the latter, a low lattice mismatched GaP/Si template can be used to circumvent the
problem of anti-phase domain formation [1].
In this contribution, we present a number of critical parameters involved in the formation by
Molecular Beam Epitaxy of a high quality 90° misfit dislocation array at the interface between GaSb
and GaAs or GaP substrates [2]. We show that, among these parameters, the initial GaAs or GaP
surface treatment is crucial for attaining a rapid and full relaxation of GaSb templates. In particular,
with AFM, TEM and RHEED investigations, we have analyzed the dislocation networks generated
using either a Ga-rich or a Sb-rich surface preparation and found that among the two conditions the
latter leads to a better accommodation (fig. 1). The influence of other growth parameters, such as
growth temperature, growth rate and V/III ratio, will be discussed in order to determine the important
mechanisms in the relaxation process. Under optimized growth conditions, AlSb/InAs heterostructures
were grown on both type of substrates exhibiting a 77K electron mobility of 108 000 cm 2 .V.s -1 and 143
000 cm 2 .V.s -1 with a sheet carrier density of 1.5x10 12 cm -2 for GaAs and GaP substrates, respectively.
Fig 1: Misfit dislocation arrays observed along [1-10] direction at the GaSb/Ga P (a) and GaSb/GaAs (b) interfaces.
[1] I. Németh, B. Kunert, W. Stolz, and K. Volz, J. Cryst. Growth 310, 1595 (2008).
[2] S.El Kazzi, L.Desplanque, C.Coinon, Y.Wang, P.Ruterana and X.Wallart, Appl. Phys. Lett. 97, 192111 (2010).
__________________________
* Contact: ludovic.desplanque@iemn.univ-lille1.fr

Mo2.6
Defects in MBE grown InAs/GaSb superlattices
on GaSb substrates
M. Walther 1* , R. Rehm 1 , J. Schmitz 1 , J. Niemasz 1 , F. Rutz 1 , A. Wörl 1 ,
L. Kirste 1 , R. Scheibner 2 , J. Ziegler 2 , and A. Danilewsky 3
1
Fraunhofer-Institut für Angewandte Festkörperphysik, Tullastraße 72, 79108 Freiburg, Germany
2 AIM Infrarot-Module GmbH, Theresienstraße 2, 74072 Heilbronn, Germany
3
Institut für Geowissenschaften, Albert-Ludwigs-Universität Freiburg, Hermann-Herder-Straße 5,
79104 Freiburg, Germany
InAs/GaSb short-period superlattices (SLs) have proven their great potential for high performance
infrared detectors. Lots of interest is currently focused on the development of short-period InAs/GaSb
SLs for mono- and bi-spectral infrared detectors between 3 – 30 µm.
InAs/GaSb short-period superlattices can be fabricated with up to 1000 periods in the intrinsic region
without revealing diffusion limited behavior [1]. This enables the fabrication of InAs/GaSb SL camera
systems with very high responsivity, comparable to state of the art CdHgTe and InSb detectors. The
material system is also ideally suited for the fabrication of dual-color mid-wavelength infrared
InAs/GaSb SL camera systems with high quantum efficiency for simultaneous and spatially coincident
detection in two different spectral channels.
An essential point for the performance of two-dimensional focal plane infrared detectors is the number
of defective pixel in the focal plane array. Sources for pixel outages are manifold and might be caused
by the dislocation in the substrate, the epitaxial growth process itself or by imperfections during the
focal plane array fabrication process. The ultimate goal is to grow defect-free epitaxial layers on a
dislocation free large area GaSb substrate. Permanent improvement of the substrate quality and the
development of techniques to monitor the substrate quality are of particular importance. To examine
the crystalline quality of 3’’ and 4’’GaSb substrates, synchrotron white beam X-ray topography
(SWBXRT) at the synchrotron light source ANKA, Karlsruhe Institute of Technology, was employed
to measure the threading dislocation density in GaSb substrates. In a comparative defect study of 3’’
GaSb and 4’’ GaSb substrates from different development generations, a significant reduction of the
dislocation density from the older to the newest substrate material has been observed, demonstrating
the improvements in bulk crystal growth technology.
The InAs/GaSb SL detector structures are grown in a multi-wafer MBE system on 3’’ GaSb substrates.
For the fabrication of single- and dual-color thermal imaging systems in the mid-wavelength infrared
region, a manufacturable growth process for high responsivity thermal imaging systems has been
developed [2]. Details on MBE growth, emphasizing on strain adjustment in thick InAs/GaSb SLs and
the reduction of MBE related defects during growth will be presented.
An automated optical characterization technique, using a Candela CS 20 inspection system for defect
characterization after MBE growth of InAs/GaSb SLs reveals a clear correlation between the defects
generated by the MBE growth process and the defective pixel maps of the detector matrix in focal
plane array detectors after hybridization with the read-out integrated circuit. Details of the
characterization technique and strategies for further reduction of defects in InAs/GaSb SLs will be
discussed.
[1] R. Rehm, M. Walther, J. Schmitz, F. Rutz, J. Fleissner, R. Scheibner and J. Ziegler. Infrared Phys. Technol. 52, (2009)
344.
[2] F. Rutz, R. Rehm, J. Schmitz, J. Fleissner, M. Walther, R. Scheibner, and J. Ziegler, Proc. SPIE 7298, (2009) 72981R-1.
__________________________
* Contact: martin.walther@iaf.fraunhofer.de

Mo2.6
Fig 1: Transmission SWBXRT topograph of different areas on a 3’’ GaSb substrate with non-uniform defect distribution and
large dislocation free areas.
Fig 2: Overlay of the optical surface characterization obtained with the Candela CS 20 system (grey scale) on high
defect density material with the defective pixel map of a dual color test FPA to demonstrate the combination of both
characterization techniques. The size of the area corresponds to about 5000 detector pixel.
Fig 3: Thermal image of a person with a candle after data fusion of both channels (3-4µm, cyan) and (4-5µm, red) in
complementary colors.

Monday Session Mo3
Nitrides

Mo3.1
Making nitrides magnetic
Alberta Bonanni
Institut für Halbleiter- und Festkörperphysik, Johannes Kepler University, Linz - Austria
We summarise our recent work on controlling and elucidating the magnetism and the
exchange interactions in GaN and related systems grown by MOVPE and doped with either Fe [1-6]
or Mn [7-9] and co-doped with donors (Si) [3,4,8] or acceptors (Mg) [3]. In particular, we show that a
significant contribution of d orbitals to the bonding leads to the self-organized aggregation of Fe
cations at the growth surface, driving the material systems to the state of condensed magnetic
semiconductor (CMS), i.e. to a semiconducting matrix with Fe-rich nanoscale chemical or
crystallographic phase separations [10,11]. The correlation between the presence of different Fe-rich
phases with peculiar and well-defined magnetic behaviour is highlighted together with ways proposed
for the realisation of a single-phase CMS. Furthermore, we demonstrate - by employing a range of
nano characterisation tools - that Mn in GaN occupies random cation positions at least up to x = 3%.
We present experimental results on the determination of the coupling strength between Mn ions in
GaN and show how by co-doping with Si the dominant interaction can be tuned from ferromagnetic to
antiferromagnetic [8].
[1] A. Bonanni, M. Kiecana, C. Simbrunner, Tian Li, M. Sawicki, M. Wegscheider, M. Quast, H.
Przybylinska, A. Navarro-Quezada, R. Jakieła, A. Wolos, W. Jantsch, and T. Dietl, Paramagnetic
GaN:Fe and ferromagnetic (Ga,Fe)N: The relationship between structural, electronic, and magnetic
properties, Phys. Rev. B 75, 125210 (2007).
[2] W. Pacuski, P. Kossacki, D. Ferrand, A. Golnik, J. Cibert, M. Wegscheider, A. Navarro-Quezada,
A. Bonanni, M. Kiecana, M. Sawicki, T. Dietl, Observation of strong-coupling effects in a diluted
magnetic semiconductor (Ga,Fe)N, Phys. Rev. Lett. 100, 037204 (2008).
[3] A. Bonanni, A. Navarro-Quezada, Tian Li, M. Wegscheider, R.T. Lechner, G. Bauer, Z. Matej, V.
Holy, M. Rovezzi, D’Acapito, M. Kiecana, M. Sawicki, and T. Dietl, Controlled aggregation of
magnetic ions in a semiconductor. Experimental demonstration, Phys. Rev. Lett. 101, 135502 (2008).
[4] M. Rovezzi, F, D’Acapito, A. Navarro-Quezada, B. Faina, T. Li, A. Bonanni, F. Filippone, A.
Amore Bonapasta, T. Dietl, Local structure of (Ga,Fe)N and (Ga,Fe)N:Si investigated by x-ray
absorption fine structure spectroscopy, Phys. Rev. B 79, 195209 (2009).
[5] A. Navarro-Quezada, W. Stefanowicz, Tian Li, B. Faina, M. Rovezzi, R. T. Lechner, T. Devillers,
F. d’Acapito, G. Bauer, M. Sawicki, T. Dietl, and A. Bonanni, Embedded magnetic phases in
(Ga,Fe)N: the key role of growth temperature, Phys. Rev. B 81, 205206 (2010).
[6] I. A. Kowalik, A. Persson, M.A. Nino, A. Navarro-Quezada, B. Faina, A. Bonanni, T. Dietl, D.
Arvanitis, Element specific characterization of heterogeneous magnetism in (Ga,Fe)N films,
arXiv:1011.0847.
[7] W. Stefanowicz, D. Sztenkiel, B. Faina, A. Grois, M. Rovezzi, T. Devillers, A. Navarro-Quezada,
T. Li, R. Jakieła, M. Sawicki, T. Dietl, and A. Bonanni, Magnetism of dilute (Ga,Mn)N, Phys. Rev. B
81, 235210 (2010).
[8] A. Bonanni, M. Sawicki, T. Devillers, W. Stefanowicz, B. Faina, Tian Li, T. E. Winkler, D.
Sztenkiel, A. Navarro-Quezada, M. Rovezzi, R. Jakiela, A. Meingast, G. Kothleitner, and T. Dietl,
Ga 1−x Mn x N – Experimental Probing of Exchange Interactions between Localized Spins in a Dilute
Magnetic Insulator, arXiv:1008.2083.
[9] J. Suffczynski, A. Grois, W. Pacuski, A. Golnik, J. A. Gaj, A. Navarro-Quezada, B. Faina, T.
Devillers, A. Bonanni, Effects of s, p -- d and s -- p exchange interactions probed by exciton
magnetospectroscopy in (Ga,Mn)N, arXiv:1011.4850, Phys. Rev. B in print
[10] A. Bonanni, Topical Review, Ferromagnetic nitride-based semiconductors doped with transition
metals and rare earths, Semicond. Sci. and Technol. 22, 41 (2007).
[11] A. Bonanni and T. Dietl, A story of high-temperature ferromagnetism in semiconductors, Chem.
Soc. Rev. 39, 528 (2010).

Mo3.2
Optical and structural properties of III-nitride nanowires and
nanowire heterostructures
F. Furtmayr 1,2,* , J. Teubert 1 , J. Arbiol 3 , P. Komninou 4 ,
M. Stutzmann 2 , and M. Eickhoff 1
1
I. Physikalisches Institut, Justus-Liebig-Universität, 35392 Giessen, Germany
2 Walter-Schottky-Institut, Technische Universität München, 85748 Garching, Germany
3 Institut de Ciència de Materials de Barcelona, CSIC, Campus de la UAB, 08193 Bellaterra, CAT, Spain
4 Aristotle University of Thessaloniki, Department of Physics, GR-54124 Thessaloniki, Greece
Group III-nitride nanowires (NWs) and nanowire heterostructures (NWHs) have attracted increasing
scientific interest due to their superior material properties compared to their two-dimensional
counterparts. In particular, they feature a substantially reduced density of structural defects, thereby
offering a promising approach for the realization of efficient light emitters and improved
nanoelectronic devices. Furthermore, they are an ideal model system for the investigation of basic
material properties.
We will discuss the nucleation process and growth details for undoped GaN NWs grown by catalystfree
plasma-assisted molecular beam epitaxy (PAMBE) on Si(111) and summarize the influence of
both Si- and Mg-doping on the NW growth, morphology, and defect formation as well as their
photoluminescence (PL) characteristics.
The formation of nanodiscs (NDs) embedded in these nanowires as well as the understanding and
control of their optical properties are basic requirements for the realization of NW-based
optoelectronic devices. We have realized GaN NDs in NWHs with Al x Ga 1-x N barriers of different Alconcentrations
x and have analyzed the carrier confinement in detail. Temperature-dependent PL
measurements were carried out for NWHs with different ND heights and different barrier
compositions and will be compared to numerical simulations based on the results of a structural
analysis by high resolution transmission electron microscopy (HRTEM) as input parameters (Fig. 1).
An increase of the Al content in the barriers from 8% to 34% increases the PL emission energy from
3.53 eV to 3.73 eV for 2 nm thick NDs. In samples with AlN barriers, the ND emission shows a red
shift even below the GaN bandgap for a ND thickness of 3.5 nm.
The emission characteristics are strongly affected by the presence of a coaxial Al(Ga)N shell which
influences the strain state and thereby the PL emission properties for samples with high Al content.
The structural and optical properties of InGaN NDs in GaN NWs will also be discussed.
a
b
20 nm
2nm
Fig 1: a) HAADF STEM image (Z-contrast) of one NW with nine 1.7 nm thick GaN nanodisks surrounded by AlN barriers.
The GaN appears bright, the surrounding AlN dark. b) HRTEM detail of a ND in the same NW. c) PL spectra of four GaN /
Al x Ga 1-x N (x = 0.14) NWs with different well thicknesses.
(c)
__________________________
* Contact: furtmayr@wsi.tum.de

Mo3.3
Specialized MBE system for growth of high quality III-N
heterostructures
A. Alexeev 1 , D. Krasovitsky 2 , S. Petrov 1* and V. Chaly 2
1 SemiTEq JSC, St.Petersburg, Engels av., 27, Russia
2 Svetlana-Rost JSC, St.Petersburg, Engels av., 27, Russia
SemiTEq JSC is the Russian leading supplier of R&D MBE Systems for the growth of different
material systems such as InAlGaN, InAlGaAs, wideband II-VI etc. This work presents the use of
SemiTEq MBE System STE3N2 specially designed for high temperature III-nitrides growth. The
unique features of this System include very high (extreme for typical MBE) range of substrate
temperatures (up to 1200 0 С defined by IR pyrometer) as well as N/III ratios. The result of use the
STE3N2 in production of power microwave transistors based on AlN/AlGaN/GaN/AlGaN DHFET
heterostructures are also shown.
One of the main problems in manufacturing of GaN-based devices up to date is the lack of low
cost, lattice matched substrates to III-nitrides. Heteroepitaxy of nitrides on mismatched substrates in
spite of using of special procedures on initial growth stage results in high dislocation density (up to
10 8 -10 10 cm -2 ) and this affects on the device quality and reliability. Moreover, typical growth
temperatures in MBE are much lower as compared with MOCVD. It leads to insufficient surface
mobility of atoms and worse coalescence of nucleation blocks on initial growth stage which, in turn,
results in increased active layer dislocation density.
We present the results of growing “thick” (more than 200 nm) high temperature AlN layers on
sapphire and SiC at 1100-1150 0 C following by appropriate sequence of transition AlGaN layers
(including superlattices) resulting in high quality final GaN layer. Dislocation density in this layer is
reduced to 1.5-2 order of magnitude in comparison with growth on the traditional thin AlN nucleation
layer. Significant reduction of dislocation density has led to substantial electron mobility increase in
active GaN layers. Maximum electron mobility in low silicon doped 1.5 mkm thick GaN layer is in the
range 600-650 сm 2 /V . s with carrier concentrations 3-5 . 10 16 cm -3 . These data correspond to the best
values reached up to date for MBE and confirms high crystal quality of a grown material.
Design and aluminum content modification of the barrier AlGaN layer in AlN/AlGaN/GaN/AlGaN
DHFET heterostructues allow to change the electron mobility and sheet concentration in two
dimensional electron gas formed at upper interface GaN/AlGaN in the range of 1300-1700 cm 2 /V . s
and 1.0-1.8 . 10 13 cm -2 , respectively, that allows to controllable change of sheet resistance in the range of
230-400 Ohm/sq.
Using these high quality DHFET heterostructues grown on SiC substrates 30 MHz-4.0 GHz
broadband amplifiers with gain 17-25 dB, P out = 2.5 W and efficiency 30% were received. These
devices have shown reliable operation during more than 3500 hours under temperature acceleration
test conditions.
__________________________
* Contact: petrov@semiteq.ru

Mo3.4
Improved luminescence and thermal stability of MBE-grown
semipolar (11-22) InGaN quantum dots
A.Das 1,* , Y. Kotsar 1 , A. Lotsari 2 , Th. Kehagias 2 , Ph. Komninou 2 , and E. Monroy 1
1
CEA-Grenoble, INAC/SP2M/NPSC, 17 rue des Martyrs, 38054 Grenoble cedex 9, France
2 Physics Department, Aristotle University of Thessaloniki, 54124 Thessaloniki, Greece
The most widely employed growth plane for GaN based optoelectronic structures is the polar
(0001) plane. The main drawback of this crystallographic orientation is the presence of internal
electric field in heterostructures, stemming from the difference in spontaneous and piezoelectric
polarization along the axis. In order to eliminate the deleterious effects of polarization,
structures grown on alternative crystal planes different from polar plane can be used. Semipolar
structures, those whose growth direction forms an angle different from 0° or 90° with the c axis, are
characterized by reduced polarization-related electric fields and better crystalline quality in
comparison with the nonpolar planes. Semipolar InGaN/GaN quantum wells (QWs) have been
demonstrated as potential active media for light emitting devices in the green/yellow spectral region.
Quantum dot (QD) structures offer a possibility to enhance the internal quantum efficiency thanks to
the three-dimensional (3D) carrier confinement which prevents carrier diffusion towards non-radiative
recombination centres. Our current work describes the characteristics of semipolar (11-22)-oriented
InGaN/GaN QDs grown by plasma-assisted MBE. For comparison purposes, polar (0001) InGaN QDs
were grown simultaneously with semipolar QDs. Two growth methods were considered, namely hightemperature
growth (substrate in the 650-590°C range) and low-temperature growth (500-450°C).
For the generation of the QDs at high temperature, the Ga flux was fixed at 30% of the
stoichiometric value and the In flux was tuned close to the stoichiometry. Therefore, the Stranski-
Krastanow transition is forced by the lattice mismatch, in spite of the slightly metal-rich atmosphere
and the well-known surfactant effect of In which promotes two-dimensional growth. For the growth of
the GaN spacer, the Ga flux was fixed at the stoichiometric value. At this high growth temperature,
InGaN decomposition is severely active, so that the In mole fraction in the QDs is drastically
influenced by the substrate temperature. HRTEM images (Fig 1) reveal the presence of 3D islands.
The polar QDs are with well-defined facets; the base diameter and height are 15-20 nm and 2-2.2 nm
respectively. The semipolar QDs are found to nucleate on flat (11-22) plane as well as at surface
depressions induced by the threading dislocations. The emission from semipolar InGaN QDs is
systematically blue-shifted in comparison to the respective polar samples grown simultaneously, as
illustrated in Fig 2. This is due to different In incorporation in the two crystallographic orientations [1].
The temperature dependence of the PL-peak energy, in the case of polar InGaN QWs and polar and
semipolar InGaN/GaN QDs, all emitting at 420 nm, is shown in Fig 3. At higher growth temperature
both polar and semipolar QD samples present similar thermal activation energy (E A ~ 80 meV),
indicating that the confinement of the electrons in the QDs is comparable, and higher than that of
InGaN/GaN QWs (E A ~ 20 meV). However, the PL thermal extinction rate is higher in the case of
semipolar QDs, which is consistent with a higher density of structural defects.
In the case of low temperature growth conditions, In desorption is negligible. The In and Ga fluxes
are hence adjusted to the desired mole fraction. Under these conditions, we have demonstrated a series
of semipolar QDs and the emission covers the full visible spectrum, from UV to IR (3.5 eV - 1.6 eV in
Fig 4). In contrast, the emission from polar QDs grown simultaneously is mostly merged with yellow
band. Semipolar QDs grown at low temperature showed an enhanced thermal stability of their
luminescence, which improves for longer emission wavelength as a result of the higher confinement
(Fig 5). These results indicate that the low temperature growth conditions are not compatible with
polar plane (0001) which provide a favorable environment to semipolar plane (11-22) to enhance the
internal quantum efficiency of InGaN/GaN nanostructures.
[1] A. Das et al., Appl. Phys. Lett., 96, 181907 (2010).
__________________________
* Contact: aparna.das@cea.fr

Mo3.4
Fig 1: HRTEM images from (a) polar InGaN QDs and (b) polar InGaN QWs, taken along [11-20] Al2O3 zone axis and (c)
semipolar sample, taken along [1-100] GaN . In (c) semipolar QDs are seen to nucleate on the (11-22) plane (left handside), as
well as on inclined planes introduced due to ascending threading dislocations.
Normalized PL Intensity
T = 7 K
Semipolar
Polar
Normalized PL Intensity
1
0.1
0.01
Polar InGaN/GaN QDs
Semipolar InGaN/GaN QDs
Polar InGaN/GaN QWs
350 400 450 500
Wavelength (nm)
Fig 2: PL spectra from semipolar and polar QDs grown at
high temperature.
1E-3
0.01 0.1
1 / Temperature (K -1 )
Fig 3: Variation of PL intensity with temperature for polar
and semipolar QDs grown at high temperature. Polar
QWs are included as a reference.
Normalized PL Intensity
Normalized PL Intensity
10 0
10 -1
10 -2
Emission Wavelength
385 nm
410 nm
515 nm
580 nm
1.2 1.6 2.0 2.4 2.8 3.2 3.6
Energy (eV)
10 -3
0.01 0.1
1 / Temperature (K -1 )
Fig 4: PL spectra from semipolar QDs grown at low
temperature.
Fig 5: Variation of PL intensity with temperature for
semipolar QDs grown at low temperature emitting at
various wavelengths.

Mo3.5
Group III-nitrides growth on N-polar substrates
C. Chèze 1,* , M. Sawicka 1,2 , M. Siekacz 1,2 , H. Turski 2 , G. Cywiński 2 , B. Grzywacz 2 ,
S. Grzanka 2 , I. Dzięcielewski 2 , B. Łucznik 1,2 , M. Boćkowski 1,2 , C. Skierbiszewski 1,2
1
TopGaN Ltd, Sokołowska 29/37, 01-142 Warszawa, Poland
2 Institute of High Pressure Physics, PAS, Sokołowska 29/37, 01-142 Warszawa, Poland
The growth of nitrogen-polar (N-polar) group III-nitrides by plasma-assisted molecular beam
epitaxy (PAMBE) is currently raising a lot of interest. One of the advantages of the growth along the
[000-1] over the [0001] direction is the highest thermal dissociation limit of N-face InN [1], that
extends as well the growth window for N-polar InGaN. A recent comparative study of InGaN grown
by PAMBE in both polarities revealed that In incorporation in N-face InGaN is indeed more efficient
than in metal-face InGaN grown at the same temperature. Moreover, the photoluminescence intensity
of the N-face samples could be as much as two orders of magnitude higher than the one of metal-face
samples [2]. Another benefit of the growth along the [000-1] direction lies in the formation of
polarization-induced three dimensional hole gas at the interface of N-polar GaN-AlGaN with graded
Al content. This method proposed by Simon et al [3] proves very efficient p-type doping. However, N-
face III-nitrides usually suffer from poor morphology and may be more prone to incorporate intrinsic
nitrogen-related defects [4].
In this work we study the growth of GaN and InGaN quantum wells on N-polar (000-1) GaN
by PAMBE. Similarly to the growth of nitrides on Ga-polar face (0001), GaN and InGaN structures
were grown in the metal-rich regime. These experiments were carried out on N-polar GaN freestanding
HVPE substrates prepared by a procedure developed at TopGaN that comprises mechanical
polishing, chemo-mechanical polishing and wet-etching. The substrate surface miscut was 0.7°
towards m-direction [10-10].
Atomic force microscopy (AFM) characterization of the samples reveals that we have
achieved high quality step-flow growth of GaN. The surface of the GaN layers grown by PAMBE on
N-polar substrates is extremely smooth (Fig. 1) and the root mean square roughness value measured
on the area of 5×5 µm 2 is 0.23 nm. We clearly observe single atomic steps flowing towards the [10-10]
direction when every second step edge is either straight or jagged. Similar observations supported by a
step-flow model are available for Ga-polar (0001) GaN [5]. Such surface morphology can be
explained by the growth anisotropy of two alternating types of step-edges characterized by their
respective number of dangling bonds [5]. In addition, the photoluminescence (PL) spectra at room
temperature of the N-polar GaN layer (Fig 2) exhibit a sharp peak at 363 nm with a full width at half
maximum value of 4.53 nm.
We will discuss the effect of the III/V ratio on the growth mode of N-polar GaN. We will also
compare the growth and characteristics of single InGaN quantum wells on N- and Ga-polar substrates.
These preliminary results open up new possibilities for more efficient N-polar based deep ultra-violet
and green optoelectronic applications.
[1] J. Simon, V. Protasenko, C. Lian, H. Xing, and D. Jena, Science 327, 60 (2010).
[2] D. Y. Nanishi, Y. Saito, T. Yamaguchi, M. Hori, F. Matsuda, T. Araki, A. Suzuki, and T. Miyajima, Phys. Stat. Sol. 200,
202 (2003).
[3] D. N. Nath, E. Gür, S. A. Ringel, and S. Rajan, Appl. Phys Lett. 97, 071903 (2010).
[4] A. R. Arehart, T. Homan, M. H. Wong, C. Poblenz, J. S. Speck, and S. A. Ringel, Appl. Phys Lett. 96, 242112 (2010).
[5] M.H. Xie, S.M. Seutter, W.K. Zhu, L.X. Zheng, H. Wu, S.Y. Tong, Phys. Rev. Lett. 82, 2749 (1999).
Acknowledgements
This work was supported partially by the Polish Ministry of Science and Higher Education Grant No IT 13426
and the European Union within European Regional Development Fund, through grant Innovative Economy
(POIG.01.01.02-00-008/08) and within FP7-PEOPLE-IAPP-2008 through grant SINOPLE 230765.
__________________________
* Contact: caroline@unipress.waw.pl

Mo3.5
Fig. 1: AFM scan of the N-polar GaN layer grown by PAMBE under Ga-rich conditions presenting parallel atomic steps
flowing towards m-direction [10-10]. Every second step edge is either straight or jagged. The scanned area is 500×500 nm 2 .
Fig. 2: PL at room temperature of the N-polar GaN layer.

Mo3.6
MBE growth of GaN using in-situ SiN treatments
F. Semond*, E. Frayssinet, M. Leroux, Y. Cordier, M. Réda Ramdani, J.C. Moreno,
S. Sergent, B. Damilano, P. Vennéguès, O. Tottereau and J. Massies
CRHEA/CNRS, rue Bernard Gregory, Sophia Antipolis, 06560 Valbonne, France
Commercially available GaN-based devices are mostly grown by MOCVD. One of the major
reasons why MOCVD is the technique of choice for GaN growth is because this growth method
allows for stronger defect reduction when GaN layers are grown on foreign substrates (Al 2 O 3 , SiC and
Si). Actually, defect reduction is mainly carried out using in-situ SiN treatments which turn out to be
very effective to promote a 3D growth mode. Even if the growth mechanism responsible for the 3D
growth mode is still controversial, it results in the bending of dislocation which forces them to interact
each other during coalescence resulting in a strong reduction of the dislocation density. As far as we
know, nobody so far succeeded to take advantage of such SiN treatments when GaN growth is
performed by MBE. Usually using MBE, as soon as a strong 3D growth mode is activated, it becomes
very difficult to recover a perfect 2D growth mode and resulting layers exhibit a poor overall quality.
In this work, using ammonia-MBE, we demonstrate for the first time that in-situ SiN
treatments could be successfully used to reduce the dislocation density and to significantly improve
the overall quality of GaN layers grown by MBE on foreign substrates. The structural, optical and
electrical properties of GaN layers grown with and without SiN treatments are discussed. When
applied to the growth of GaN on silicon substrates, the reduction of the dislocation density
significantly influences the strain state of GaN layers indicating a strong relationship in between
dislocation density and strain relaxation. Also, from a more fundamental point of view, the fact that
SiN treatments as well as the resulting 3D growth mode could be followed in-situ by RHEED during
growth could bring new insights regarding this SiN treatment procedure which is massively used by
MOCVD growers. For instance, we discovered that the in-situ SiN treatment carried out by MBE
leads to the formation of a 3x1 surface reconstruction which turns out to be stable under air exposure.
Surprisingly, a similar 3x1 surface reconstruction is observed on MOCVD GaN layers ended by a
standard in-situ SiN treatment carried out by MOCVD and then transferred into an MBE reactor.
These observations indicate first, that SiN treatments carried out by MOCVD and MBE lead to a
reconstructed surface and secondly, the reconstruction is very similar in terms of structure and
composition whatever the growth technique used. Of course density and size of 3D GaN islands
formed by MBE and by MOCVD are quite different, due to large differences in terms of surface
diffusion lengths, but in both cases coalescence could be achieved.
This result brings a new tool for MBE growers who are interested to decrease the dislocation
density in their GaN layers grown on foreign substrates. Also, it is anticipated that the understanding
of the mechanisms responsible for the strong 3D growth and the defect reduction during coalescence is
potentially of interest for many kinds of highly mismatched hetero-epitaxial growth.
__________________________
* Contact: fs@crhea.cnrs.fr

Quadrupole Mass Spectrometers for MBE
Hiden Analytical manufacture a range of high sensitivity, high stability Quadrupole Mass
Spectrometers configured specifically for MBE environments. The XBS series Flux Monitor and
HALO series Residual Gas Analysers represent the cutting edge in MBE process control and
vacuum diagnostics and are available to both end user and OEM customers alike.
XBS Beam Flux Monitor
The XBS series monitors for MBE are quadrupole mass spectrometers featuring an applicationspecific
ionization source with wide acceptance angle for simultaneous monitoring and control
of multiple source beams.
Beam acceptance is through a wide 70 o cone,
an interchangeable 3-dimensional pepperpot
interface being factory-configured to exactly
match the source positions for each process
chamber. An integral water-cooled shroud
enables probe mounting close to the source
positions and/or close to the heater substrate
stage. Chamber insertion length is
configurable to up to 550 mm.
The system is fully UHV compatible and offers a mass range to 320 amu, 10-decade dynamic
range and PC control together with up to 16 analogue output channels for beam control
reference.
HALO Residual Gas Analyser
The HALO series MBE residual gas analysers feature custom ion source shrouds and MBE
compatible construction for contamination resistance and increased lifespan.
They are available in 100,
200 and 300 amu mass
range options and offer
significantly improved
performance when
compared with standard,
unshielded RGA’s.
Hiden’s Windows
MASsoft control software
is supplied as standard
and includes fast data
acquisition through either user configured acquisition files or pre-set modes selected by icon.
420 Europa Blvd., Warrington WA5 7UN, England
Tel. +44 (0)1925 445225 Fax. +44 (0)1925 416518
Email info@hiden.co.uk Web www.HidenAnalytical.com

Tuesday Session Tu1
Arsenides II

Tu1.1
Advances of dilute-nitrides MBE technology and related device
applications
M. Guina * , V.-M. Korpijärvi, J. Puustinen, A. Aho, T. Leinonen, and A. Tukiainen
Optoelectronics Research Centre, Tampere University of Technology,
Korkekoulunkatu 3, Tampere 33720, Finland.
Dilute-nitride (GaInNAs/GaAs) heterostructures are providing unexpected opportunities for the
development of novel optoelectronics devices outside the telecomm lasers arena, which, so far, has
been the main driving force for studying this material system. For example, dilute-nitrides have
recently enabled the development of reliable vertical-external-cavity surface-emitting lasers
(VECSELs) generating multi-watt single-mode radiation at 1180-1220 nm [1, 2]. Low N content (N <
1.5 %) dilute-nitrides can provide a practical approach for developing lasers, amplifiers, and
electroabsorption modulators required for optical integration; ultra-short interconnections based on
passive silicon waveguides could operate at wavelengths below 1250 nm rendering possible the use of
high quality GaInNAs/GaAs heterostructures containing a small amount of N. We should also note
here that a GaAs-based optical integration platform could provide important means for cost reduction
owing to compatibility with 6 inch microelectronics technology and CMOS electronics. Quite
remarkably, even the detrimental effects the N incorporation has on crystal quality can be exploited
positively in ultrafast optics applications. Thus, dilute-nitrides can be used to tailor the recovery time
of semiconductor saturable absorber mirrors (SESAMs) operating in a broad wavelength range from
1.1 µm to 1.6 µm. Finally, we should mention that dilute-nitrides continue to hold an important
potential for the development of ultra-high efficiency multi-junction solar cells [3]. This high-impact
application requires sustained efforts to improve the optical and electrical properties of dilute-nitrides
combined with deployment of advanced device concepts, such as quantum-well absorbing regions.
MBE is essential for all these developments owing to its unique ability to control and monitor the
growth conditions and the N incorporation.
The presentation reviews recent advances concerning the understanding of nitrogen incorporation
mechanisms, and the physics of dilute nitride heterostructures and related devices. In particular, we
discuss the optimization of plasma assisted MBE processes used to fabricate high-power 1180 nm
VECSELs, 1030-1300 nm SESAMs with tailored recovery time, lattice-matched solar cells absorbing
the 1 eV solar spectrum, and ultrafast laser diodes for optical interconnects.
[1] M. Guina, T. Leinonen, A. Härkönen, and M. Pessa, New J. Phys. 11, 125019 (2009)
[2] V.-M. Korpijärvi, T. Leinonen, J. Puustinen, A. Härkönen, and M. Guina, Opt. Exp. 24. 25633 (2010)
[3] D. J. Friedman, J. F. Geisz, S. R. Kurtz, and J. M. Olson, J. Crystal Growth 195, 409 (1998).
__________________________
* Contact: mircea.guina@tut.fi

Tu1.2
Current-injection lasing in GaAs quantum dots grown by droplet
epitaxy
M. Jo * , T. Mano and K. Sakoda
National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047,Japan
Quantum dot (QD) lasers have many advantages compared to existing solid-state lasers, such as
high-speed modulation, low-threshold current, and high-temperature characteristics. GaAs/AlGaAs
QDs emitting in the wavelength range of 600-800 nm are an attractive candidate for QD lasers for
on-chip optical interconnect applications. To achieve the lasing, uniformly sized QDs with high density
and high quality are desirable. So far, photopumped laser action was reported for GaAs quantum rings
on (001) substrates 1 and GaAs QDs on (311)A substrates. 2 The former improved the size uniformity
using low As pressure, while the latter realized high density of 10 11 cm -2 with good uniformity on (311)A.
However, no electrically pumped lasing has been observed yet.
Here, we report electrically pumped laser action of self-assembled GaAs/AlGaAs QDs grown on
(001) substrates by droplet epitaxy. High-quality GaAs QDs were fabricated by thin AlGaAs capping
and subsequent in-situ annealing. 3 High-temperature annealing also flattened the top of the QDs,
resulting in the formation of height-controlled QDs. Furthermore, an artificial GaAs wetting layer of 2
nm was introduced to effectively enhance the dot volume as well as relieve the dot height fluctuations.
All these improved growth techniques lead to electrically injected lasing in GaAs QDs.
Fig 1: 500×500-nm AFM image of uncapped GaAs QDs.
The dot density was 2×10 10 cm -2 .
Fig 2: Lasing characteristics of GaAs/AlGaAs QDs at 77
K. The sample contains five-fold stacked QDs in the active
core. The inset shows L-I characteristics.
[1] T. Mano, T. Kuroda, M. Yamagiwa, G. Kido, K. Sakoda, and N. Koguchi, Appl. Phys. Lett., 89, 183102 (2006).
[2] T. Mano, T. Kuroda, K. Mitsuishi, Y. Nakayama, T. Noda, and K. Sakoda, Appl. Phys. Lett., 93, 203110 (2008).
[3] M. Jo, T. Mano, and K. Sakoda, J. Appl. Phys., 108, 083505 (2010).
__________________________
* Contact: JO.Masafumi@nims.go.jp

Tu1.3
MBE growth of high power Modelocked Integrated External-
Cavity Surface Emitting Laser (MIXSEL) with 6.4 W output power
M. Golling * , B. Rudin, V. J. Wittwer, D. J. H. C. Maas, Y. Barbarin,
M. Hoffmann, O. D. Sieber, T. Südmeyer, U. Keller
Department of Physics, Ultrafast Laser Physics Lab, ETH Zurich, Switzerland
Semiconductor disk lasers (also called VECSELs) combine the benefits of diode-pumped solid-state
lasers and semiconductor technologies, resulting in wavelength flexibility and high power operation
with excellent beam quality. However, the laser contains two separate semiconductor elements in a
folded cavity, which is a challenge for cost-efficient high volume fabrication, as well as for reaching
high repetition rates. Modelocked integrated external-cavity surface emitting lasers (MIXSELs) [1]
combine gain and absorber in one semiconductor structure, enabling modelocking in a simple straight
cavity and the possibility of a quasi-monolithic design.
For the MIXSEL, the beam diameters on gain and absorber are the same. For stable modelocking, the
absorber has to saturate faster than the gain. In the first MIXSEL this was solved by placing the
quantum dot (QD) saturable absorber layer in a resonant field enhancement. Unfortunately, the
resonant design leads to a large sensitivity towards growth errors.
Here we present the MBE growth of a substantially improved MIXSEL. We developed QD saturable
absorbers with lower saturation energy, which allow for an antiresonant design. This relaxes the
demands on the growth accuracy and avoids narrow resonances in the dispersion.
We also improved the processing using wafer removal and mounting of the 8-µm thick
MIXSEL structure directly onto a CVD-diamond heat spreader. The simple straight cavity
with only two components has generated 28-ps pulses at 2.5-GHz repetition rate and an
average output power of 6.4 W, which is higher than for any other modelocked semiconductor
laser.
MIXSEL concept
laser
absorber region
single InAs quantum dot layer
gain region
7 In 13 Ga 87 As quantum wells
50 mm
pump
output coupler
laser DBR
QD abs.
pump DBR
QW gain
AR section
AlAs Al 20
Ga 80
As GaAs
AlAs Al 20
Ga 80
As GaAs
MIXSEL chip
SEM image of MIXSEL semiconductor structure
heat sink
pump
output coupler
1 µm
laser DBR
absorber pump DBR
gain
AR section
MIXSEL chip
a b
Fig 1: (a) MIXSEL concept and design: The MIXSEL semiconductor structure contains two highly reflecting
distributed Bragg reflectors (DBRs), a quantum dot (QD) saturable absorber layer, a quantum well (QW) gain section
and an anti-reflective (AR) coating. The laser DBR reflects the laser light and forms the laser cavity together with the
external output coupler. The pump DBR is placed between the QD saturable absorber and the QW gain layers to
prevent bleaching of the saturable absorber by the pump light. The white oscillations in the upper sketches represent
the square of the electric field of the laser light. The QD layer as well as the QWs are placed in antinodes.
(b) MIXSEL cavity: The simple linear cavity is formed by the MIXSEL chip and an external output coupler. The
MIXSEL structure is optically pumped under an angle of 45°.
[1] D. J. H. C. Maas, A.-R. Bellancourt, B. Rudin, M. Golling, H. J. Unold, T. Südmeyer, and U. Keller, Appl. Phys. B 88, 493-497 (2007).
__________________________
* Contact: golling@phys.ethz.ch

Tu1.4
1.55 µm lasers based on shape-engineered InAs/InAlGaAs/InP
(100) quantum dots
C. Gilfert * , V. Ivanov and J.P. Reithmaier
Technische Physik, Institute of Nanostructure Technologies and Analytics, University of Kassel, 34132 Kassel,
Germany
Very recently InP-based quantum dots (QDs) material was realized with significantly reduced size
distribution and more circular shape in-plane [1]. In this paper the first realization of QD lasers are
reported using this new type of shape-engineered InAs/InAlGaAs/InP QD material.
QD material with enhanced spectral gain is very promising for high-speed low-chirp directly
modulated diode lasers or optical amplifiers due to its higher differential gain and low linewidth
enhancement factor, respectively. All structures are grown by solid-source molecular beam epitaxy
(MBE) using valved cracking cells for generation of As 2 and P 2 . The laser is based on a separate
confinement heterostructure design. 180 nm thick In 0.528 Al 0.238 Ga 0.234 As layers on each side of the
active region serve as waveguide. The lower cladding is formed by the n-doped InP substrate, a 200
nm InP buffer and 200 nm In 0.523 Al 0.477 As layer. The upper cladding is composed of 200 nm
In 0.523 Al 0.477 As followed by 1.7 µm InP layer. 200 nm highly-doped In 0.532 Ga 0.468 As contact layer is
deposited. The active region and the inner 80 nm of the waveguide are undoped. The active region is
composed of 6 layers of InAs QDs having a nominal thickness of 4.5 ML each. They are separated by
20 nm In 0.528 Al 0.238 Ga 0.234 As. Pulsed characterization of broad-area devices with stripe widths of 100
µm shows good lasing properties with an emission wavelength of ~ 1.57 µm as depicted in fig. 1.
Fig. 1: Plots of the device`s length-dependent pulsed-current characteristics for evaluation of the internal parameters. Inset in
(a) depicts the emission spectrum above threshold for a 0.8 mm long cavity.
The internal characteristics of the devices feature a relatively low internal absorption of 8 cm -1, and a
transparency current density of about 300 A/cm 2 . In particular, a relatively high modal gain of 60 cm -1
was obtained, which is nearly a factor of two higher than for other reported QD laser structures [2].
This may be attributed to the improved spectral gain stemming on the lower line width of the QD
material as compared to typical QDash structures. The financial support by the EU projects GOSPEL
and DeLight is acknowledged.
[1] C. Gilfert, E.-M. Pavelescu, J.P. Reithmaier, Appl. Phys. Lett., 96, 191903 (2010)
[2] K. Merghem, A. Akrout, A. Martinez, G. Aubin, A. Ramdane, F. Lelarge, G.H. Duan, Appl. Phys. Lett. 94, 021107 (2009).
__________________________
* Contact: gilfert@ina.uni-kassel.de

Tu1.5
Broadband emission and mode-locking using controlled
distributions of InGaAs quantum dots
M.Hopkinson 1,* , M. Hugues 1 , P.D.L. Greenwood 1 , M.Krakowski 2 , M. Calligaro 2 ,
S.Bruer 3 , , M. Rossetti 4 , W. Elsäßer 3 , I. Montrosset 4
1
Department of Electronic and Electrical Engineering, University of Sheffield, Sheffield S1 3JD, UK.
2 Alcatel Thales III-V Laboratory, 91767 Palaiseau, FR
3 Technische Universität Darmstadt, Institute of Applied Physics, Schloßgartenstraße 7, 64289 Darmstadt, DE
4 Politecnico di Torino, Corso Duca degli Abruzzi 24, 10129 Torino, IT
Semiconductor quantum dots (QDs) offer a wealth of new optical and electronic phenomenon based on
strong carrier localization and zero-dimensional atomic-like density of states. One characteristic of selforganised
epitaxial QDs is an inhomogeneous distribution of dot sizes, resulting in broad emission from the
QD ensemble 1 . For those applications requiring access to a well defined QD energy this presents a major
problem. However there are a range of devices where a broadband emission spectrum is a positive
advantage. In this paper we discuss the application of broad ensembles of QDs in superluminescent diodes 2
and mode-locked lasers 3 , areas of QD photonics that have received much interest in recent years.
InGaAs quantum dot devices have been produced by Molecular Beam Epitaxy, using the typical Stranski-
Krastanow mode. A dot-in-well structure (GaAs or In x Ga 1-x As QWs) is used to control the emission energy
of these structures to cover the wavelength range from 0.95 to 1.35µm. The active regions are contained
within standard doped GaAs/AlGaAs waveguide structures. Typical QD ensembles have a
photoluminescent linewidth of 30-40meV arising from the inhomogeneous distribution of dot sizes. To
increase this linewidth further we employ two methods. In the first, we exploit gain saturation of the QD
ground state, which results in a broad emission characterised by a mixture of ground and excited states.
Using this approach, bandwidths of >100meV are possible, but the gain profile is composed of overlapping
peaks and is highly discontinuous. To smooth this highly peaked emission, we introduce a variation to the
QD energy in different layers of the QD stack. A number of methods to do this in a controlled way have
been examined and will be discussed. When applied optimally, a broad emission peak with an
approximation to a Gaussian profile may be produced 4 . This type of structure has been used successfully to
produce broadband superluminescent diodes with >100nm bandwidth and power output of several mW,
suitable for use in optical coherence tomography.
The broad gain profile of QD ensembles also lends itself naturally to the development of ultrafast sources
via the time-bandwidth relationship. Compact ultrafast semiconductor mode-locked lasers operating in the
sub-ps range are highly attractive for a range of applications from chemical sensing through to optical
communications. For passive mode locking, the device is divided into gain and absorber sections (see fig.
5) and the absorber driven with zero or reverse bias. Saturation of the gain occurs easily in QD material and
this, in combination with fast recombination dynamics results in a train of pulses of width down to 3ps
being developed 5 . The unusual nature of the broadened QD distribution leads to an unusual evolution of
this mode locking in which ES emission first dominates and with increasing injection switches to ground
state lasing with a broad region of co-existence in between. The behaviour is entirely opposite to that
normally expected, but can be simply modelled if we assume that a photo-pumping process exists between
QDs. The results suggest that a well designed and grown broadband distribution can be used to produce
highly customised mode-locking characteristics.
Support from the EU Framework & programs ‘Nano ultrabright sources’ and ‘Fastdot’ is gratefully
acknowledged.
[1] D. Leonard, M. Krishnamurthy, C. M. Reaves, S. P. Denbaars, and P. M.Petroff. Appl. Phys. Lett. 63, 3203(1993)
[2] M.Rossetti. IEEE Photonic Tech. Letts.18 1946 (2006)
[3] E. U. Rafailov et al. Appl.Phys.Letts. 87, 81107, (2005)
[4] P D. L. Greenwood et al. IEEE J. Selected topics in Quan.Electron. 16, 4 1015 (2010)
[5] S.Bruer. et al., Proc. SPIE 7720, 772011 (2010)
__________________________
* Contact: m.hopkinson@sheffield.ac.uk

Tu1.5
InGaAs
QW
7nm
1nm
InGaAs
QW
1nm
7nm
Fig 1: Typical InGaAs dot-in-well structure,
shown in cross sectional TEM
Fig 2: Method of broadening the QD emission by moving
the position of the QD within the surrounding quantum well
PL intensity (arb. unit.)
6
4
2
0
1100
1200 1300
Wavelength (nm)
Fig 3: Shift in QD emission of over 150nm in different
samples achieved by changing the position of the QD
Intensity (arb. units)
950 1000 1050 1100 1150
Wavelength (nm)
Fig 4: Broadband superluminescent diode spectrum with
increasing current injection. Device designed for 1050nm
centre wavelength suitable for ocular OCT
Fig 5: Passive mode-locked QD laser device, with
separate gain and absorber section (Fabrication- III-V lab)
Fig 6: Evolution of mode-locked laser spectrum, showing
tunable emission using an external feedback resistor in a
self-electro optic device (SEED) configuration

Tuesday Session Tu2
New trends in MBE

Tu2.1
Current developments in MBE growth of highly mismatched
materials
J.B. Rodriguez * , L. Cerutti, J.R. Reboul and E. Tournié
Institut d’Electronique du Sud (IES), Université Montpellier 2, UMR CNRS 5214, 34095 Montpellier, France
Many high-end devices rely on the epitaxy of semiconductor compound with a low lattice
mismatch both between the different epi-layers and with the substrate. This usually ensures the final
material quality and a very low defect density, but prevents the realization of a large panel of
heterostructures, or the epitaxy on highly-mismatched substrates.
In particular, the fabrication of III-V semiconductors on silicon substrate usually suffers from a very
large lattice mismatch, resulting in threading dislocation densities not compatible with the fabrication
of high performance devices. For that reason, the early studies failed to demonstrate a good quality III-
V material grown on silicon.
Lately, however, many research groups and companies have shown great interest on that topic.
On one hand, applications have multiplied, from optical interconnect to low-consumption/high speed
transistors. On the other hand, several new epitaxy techniques have made impressive progress.
The presentation will review these different techniques and will focus more particularly on the
epitaxy of antimonide-based semiconductors on silicon substrates. It has been found, indeed, that these
materials can relax the strain due to the lattice-mismatch (~13% with silicon) almost entirely at the
interface with the substrate by the formation of Lomer-type dislocations. These dislocations propagate
horizontally at the interface in a self-arranged manner, leading to a significantly reduced threading
dislocation density. The Figure 1 shows the lattice constant (as measured by RHEED) evolution with
time after the growth of AlSb on silicon has started.
Using this particular relaxation mode, we have already demonstrated that it is possible to fabricate
edge emitting lasers operating in pulsed regime at 1.55 1 and 2.3 2 µm, with a buffer layer thickness of
only 1 µm. More recently, thanks to different improvements, we managed to obtain continuous-wave
operating lasers up to 35°C (Figure 2) at 2 µm. We will discuss the progress and perspectives for this
exciting new epitaxy technique.
Lattice constant (Å)
6.2
6.1
6.0
5.9
5.8
5.7
5.6
5.5
5.4
AlSb on
1 ML
26 28 30 32 34 36 38 40
Time (s)
(a)
14
12
10
8
6
4
2
0
Relative variation (%)
Voltage (V)
6
5
4
3
2
1
1.96 1.98 2.00
Wavelength (µm)
10°C
15°C
20°C
25°C
30°C
35°C
0
0 50 100 150 200 250 300
Current (mA)
6µm x 1000µm
contact Top-top
CW
Fig 1: Evolution of the lattice constant measured by RHEED for the growth of AlSb on silicon (a), P-I-V
characteristics of Sb-based lasers grown on silicon wafers in the 10-35°C temperature range and emission spectra (b).
285mA
25°C
(b)
10
8
6
4
2
0
Power (mW/facet)
[1] L. Cerutti, J. B. Rodriguez, and E. Tournie, Photonics Technology Letters, vol. 22, no. 8, April 15, (2010),
[2] J. B. Rodriguez, L. Cerutti, P. Grech, and E. Tournié, Applied Physics Letters, 94, 061124 (2009).
__________________________
* Contact: rodriguez@ies.univ-montp2.fr

Tu2.2
MBE of semiconducting oxides
O. Bierwagen 1,4,* , T. Nagata 1,2 , M.E. White 1 , M.Y. Tsai 3 , and J.S. Speck 1
1
Materials, University of California, Santa Barbara, CA 93106, USA
2
National Institute for Materials Science, 1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan
3
Department of Electrical and Computer Engineering, University of California, Santa Barbara, CA 93106, USA
4 Paul-Drude-Institut, Hausvogteiplatz 5-7, 10117 Berlin, Germany
In recent years transparent, semiconducting oxides, such as ZnO, TiO 2 , SrTiO 3 , SnO 2, In 2 O 3 , Ga 2 O 3 ,
have received an increasing amount of scientific attention. Traditionally, many of these oxides are used
either highly doped as transparent contacts in light emitters, solar cells, displays, smart windows, or
undoped in chemical sensors. These applications, however, are based on low quality material. When
grown to a high quality, in terms of structure and impurities, transparent oxides can become true
transparent semiconductors in their own right. A high material quality allows to measure intrinsic
physics of the material (and not of defects and impurities), which helps to understand the success in
existing applications, and can spawn new applications e.g. in high performance transparent electronics.
From a physics point of view, transparent semiconducting oxides are rich with wide band-gaps, high
exciton binding energy, dipole forbidden direct transitions, unintentional n-type conductivity, issues in
p-type doping the material, surface electron accumulation layers, polaronic effects, Mott transition.
Molecular beam exitaxy is an excellent tool to synthesize high quality semiconducting oxide thin
films. The highest quality films are obtained by homoepitaxy on available bulk oxide substrates (ZnO,
TiO 2 , SrTiO 3 , and Ga 2 O 3 ). During MBE growth, the metal is typically evaporated from a standard
effusion cell, whereas oxygen is supplied through a plasma source that “activates” molecular oxgen by
cracking it into monotatomic oxygen or creating excited oxygen atoms or molecules that react with the
metal ad-atoms on the substrate.
Semiconducting oxides have a wide range of thrilling MBE growth issues such as the parasitic
formation of volatile sub-oxides [4,5], facetting [1], nucleation issues [2], low vapor pressure source
material [6], and extremely narrow growth windows for stoichiometric films [7]. This talk will
highlight these growth issues and show solutions for them, focusing on the MBE growth of In 2 O 3
[1,2,3], and touching on the MBE growth of SnO 2 [4], Ga 2 O 3 [5], TiO 2 [6], and SrTiO 3 [7,8]. Using the
example of SnO 2 , the control of the conductivity from semi-insulating to highly conductive by doping
is demonstrated [9,10], and the effect of surface accumulation layer on contacts is shown [11,12].
[1] O. Bierwagen, M.E. White, M.Y. Tsai, and J.S. Speck, Appl. Phys. Lett. 95, 262105 (2009).
[2] O. Bierwagen and J. S. Speck, J. Appl. Phys. 107, 113519 (2010).
[3] O. Bierwagen and J. S. Speck, Appl. Phys. Lett. 97, 072103 (2010).
[4] M.Y. Tsai, M.E. White, and J.S. Speck, J. Appl. Phys. 106, 024911 (2009).
[5] M.Y. Tsai, O. Bierwagen, M.E. White, and J.S. Speck, J. Vac. Sci. Technol. A 28, 354 (2010).
[6] B. Jalan, R. Engel-Herbert, J. Cagnon, and S. Stemmer, J. Vac. Sci. Technol. A 27, 230 (2009).
[7] B. Jalan, P. Moetakef, and S. Stemmer, Appl. Phys. Lett. 95, 032906 (2009).
[8] J. Son, P. Moetakef, B. Jalan, O. Bierwagen, N.J. Wright, R. Engel-Herbert, and S. Stemmer, Nature. Mater. 9, 482
(2010).
[9] M.E. White, O. Bierwagen, M.Y. Tsai, and J.S. Speck, J. Appl. Phys. 106, 093704 (2009).
[10] M.E. White, O. Bierwagen, M.Y. Tsai, and J.S. Speck, Appl. Phys. Express 3, 051101 (2010).
[11] O. Bierwagen, M.E. White, M.Y. Tsai, T. Nagata, and J.S. Speck, Appl. Phys. Express 2, 106502 (2009).
[12] T. Nagata, O. Bierwagen, M.E. White, M.Y. Tsai, and J.S. Speck, J. Appl. Phys. 107, 033707 (2010).
__________________________
*
Contact: bierwagen@pdi-berlin.de

Tu2.3
MBE growth of the topological insulator Bi 2 Te 3 on Si (111)
substrates
G. Mussler 1,* , J. Krumrain 1 , L. Plucinski 2 , and D. Grützmacher 1
1
Institute of Bio- and Nanosystems 1, Research Center Jülich, 52428 Jülich, Germany
2 Institut für Festkörperforschung, Research Center Jülich, 52428 Jülich, Germany
Recently a new state of matter called topological insulator (TI) has been theoretically predicted and
experimentally observed in a number of materials [1]. Topological insulators are characterized by
gapless surface states that show a linear energy dispersion, similar to relativistic particles. Hence,
carriers at the surface of topological insulators have unparalleled properties, such as extremely high
mobilities, or dissipationless spin-locked transport, and consequently these features may lead to new
applications in the field of spintronics or quantum computing. Concerning the Bi 2 Te 3 material system,
this narrow gap semiconductor has been traditionally investigated as a thermoelectric material.
However, very recently a TI behavior has been observed at the surface of Bi 2 Te 3 [2]. To date, the
Bi 2 Te 3 material used to study TI behavior have mainly been carried out in the form of bulk crystals
realized by means of the melt-growth or self-flux method [2,3], which results in heavily n-type doped
material due to the formation of defects. To compensate the n-type doping, the Bi 2 Te 3 material is
usually heavily doped by Sn or Ca, which strongly degrades transport properties. In order to
investigate TI properties of Bi 2 Te 3 , it is therefore desirable to grow intrinsic thin films of Bi 2 Te 3 .
Besides studying fundamental properties of TI Bi 2 Te 3 , the realization of high quality Bi 2 Te 3 epilayers
on low-cost substrates, such as silicon, is highly beneficial for device applications.
Here we will present experimental results of Bi 2 Te3 thin films grown by molecular-beam epitaxy
(MBE) onto Si (111) substrates. By a careful optimization of the growth parameters, we were able to
realize high quality single-crystal Bi 2 Te 3 epilayers. Figure 1a depicts a symmetric XRD curve of a 20
nm thin Bi 2 Te 3 film on a Si (111) substrate. Besides the XRD peaks due to the Si substrate, numerous
narrow XRD peaks originating from the Bi 2 Te 3 epilayer are observed, indicating a (001)-oriented
single crystal Bi 2 Te 3 films commensurately grown on a Si(111) substrate. Atomic force microscopy
measurements were carried out, as depicted in figure 1b. The AFM image shows atomic steps with
step heights of 1.017 nm (see figure 1c), which represents the thickness of a single Bi 2 Te 3 quintuple
layer. Most importantly, figure 1d illustrates ARPES measurements of the Bi 2 Te 3 surface. A clear
linear energy dispersion is seen, evidencing the TI behavior of the MBE-grown Bi 2 Te 3 film. In
addition, we will also present transport measurements in dependence of the applied bias. The results
indicate surface carrier transport with enhanced mobilities, which are attributed to the linear energy
dispersion at the surface.
.
[1] J. E. Moore, Nature 464, 194 (2010)
[2] D. Hsieh, D. Qian, L. Wray, Y. Xia, Y. S. Hor, R. J. Cava, and M. Z. Hasan, Nature 452 970 (2008)
[3] Y. L. Chen, J. G. Analytis, J.-H. Chu, Z. K. Liu, S.-K. Mo, X. L. Qi, H. J. Zhang, D. H. Lu, X. Dai, Z. Fang, S. C. Zhang,
I. R. Fisher, Z. Hussain, and Z.-X. Shen, Science 325, 178 (2009)
__________________________
* Contact: g.mussler@fz-juelich.de

Tu2.3
Figure 1: a) symmetric XRD curve of a Bi 2 Te 3 epilayer grown on a Si (111) substrate, b) atomic force micrograph of the
Bi 2 Te 3 surface showing atomic steps, c) height profile of these atomic steps, indicating a constant step height of 10.17 A,
which represents the height of a single Bi 2 Te 3 quintuple layer, d) ARPES measurements revealing the linear energy dispersion
behavior at the Bi 2 Te 3 surface.

Tu2.4
Atomic-scale mapping of quantum dots using direct x-ray methods
R. Clarke 1,* , D.P. Kumah 1,+ , R.S. Goldman 2 , V. Dasika 2 , C. Schlepütz 1 , Y. Yacoby 3 ,
E. Cohen 4 , and Y. Paltiel 4
1
University of Michigan, Department of Physics, Ann Arbor, MI 48109, USA
2 University of Michigan, Department of Materials Science and Engineering, Ann Arbor, MI 48109, USA
3
Hebrew University, Racah Institute of Physics, Jerusalem 91904, Israel
4
Hebrew University, Applied Physics Department, Jerusalem 91904, Israel
To carefully control the growth of quantum dots and more precisely predict their optoelectronic
behavior, a better understanding of their final structural states is required [1]. Furthermore, the effects
of strain and growth conditions can lead to an exchange of material between the deposited epilayer and
the surrounding material (i.e. capping layer and underlying substrate), resulting in modifications of the
dot composition and that of the surrounding material. The tendency of some of the constituent
elements in III-V QD systems, such as indium and antimony, to preferentially segregate also
contributes to the complex structure of these systems [2].
In this presentation, we report groundbreaking results on the determination of the internal structure of
self-assembled III-V quantum dots grown by both the Stranski-Krastanow (S-K) and droplet
heteroepitaxy (DHE) methods. For the first time, we are able to map the three-dimensional atomic
structure of the quantum dots and the interface layers directly underneath the dots, thereby revealing
their structure, chemical compostion and strain with sub-Angstrom resolution. This is accomplished by
a powerful x-ray phase retrieval technique in which x-ray diffraction intensities are measured along
substrate-defined Bragg rods and converted into three-dimensional real-space electron density maps
[3].
A comparison of the maps obtained close to the x-ray absorption edges of some of the constituent
elements permits a direct determination of the chemical profile of the system in addition to the
structure. The layer-by-layer composition and atomic structure of the dots and the dot-substrate
interface are elucidated from the 3D electron density maps obtained close to the Ga and As x-ray
absorption edges. In the classic InAs/GaAs system, formed by the S-K approach, we find that the dots
have an InAs composition, as expected, and in-plane atomic registry with the underlying substrate.
However, contrary to conventional notions of the S-K growth process, no continuous epitaxial InAs
wetting layer is observed between the dots and the GaAs substrate for high dot coverages.
A change in the stacking sequence of the atomic planes in the vicinity of the dot-substrate interface is
also observed analogous to our previously published results on InSb quantum dots grown on GaAs by
the DHE method [2]. An analysis of the electron density maps shows that, while the dots are in
excellent in-plane atomic registry with the GaAs substrate, for each atomic plane in the dots, there is
variation in the vertical atomic positions. This observation is explained in terms of a structural model
of the dots displaying curved atomic planes.
The results reported here provide new insights on the structure and morphology of quantum dots and
reveal subtle atomic-level differences between dots formed by the S-K and DHE methods. The
findings also shed new light on the role of the wetting layer in the formation of quantum dots.
[1] J. Stangl, V. Holy, and G. Bauer, Rev. Mod. Phys., 76 725 (2004).
[2] D. P. Kumah, S. Shusterman, Y. Paltiel, Y. Yacoby, R. Clarke, Nature Nanotechnology 4 835 (2009).
[3] Y. Yacoby, M. Sowwan, E. Stern, J. Cross, D. Brewe, R. Pindak, J. Pitney, E. Dufresne, R. Clarke, Nature Materials 1,
99 (2002).
__________________________
* Contact: royc@umich.edu
+ Present address: Yale University, Department of Applied Physics, New Haven, CT 06520, USA

Tu2.4
Fig. 1: Cross-section Scanning Tunneling
Micrograph (XSTM) of stacked S-K quantum
dots. Scale: 25nm x 30 nm.
FIG. 2. Illustration of the substrate (GaAs)-QD (InAs) structure: (a) [010] cut through the x-ray determined
electron density map showing G-III atomic positions. (b) Vertical electron density line profile
along the G-III atomic positions indicated by the red and blue lines in (a). The asterisks denote incomplete
substrate surface layers. The inset shows the layer positions relative to bulk GaAs determined from x-ray (red)
and XSTM(black). (c) A model of the substrate/QD atomic planes determined from (a) showing the bowed
atomic planes within the dot and the displacement of the dot layers by 0.5 UC below the GaAs surface. Note that
the horizontal scale in (c) is compressed by a factor of 2 relative to the vertical scale.

Tu2.5
II-VI-based microcavities for the blue-violet spectral range
S. Klembt 1* , C. Kruse 1 , M. Seyfried 2 , K. Sebald 2 , J. Gutowski 2 and D. Hommel 1
1 Institute of Solid State Physics, Semiconductor Epitaxy, University of Bremen,
Otto-Hahn-Allee NW1, 28359 Bremen, Germany
2 Institute of Solid State Physics, Semiconductor Optics, University of Bremen,
Otto-Hahn-Allee NW1, 28359 Bremen, Germany
The aim is to realize high-quality microcavities containing ZnSe quantum wells (QWs) as the active
region. The main objectives are the observation of polariton condensation effects at elevated
temperatures and the realization of blue vertical-cavity surface-emitting lasers (VCSELs). The samples
are grown by molecular beam epitaxy (MBE).
The ZnSe material system is known to be well suited for strong-coupling experiments due to the large
exciton binding energy E b
> 25 meV and a Rabi splitting per QW of 10 meV. In this context a new
distributed Bragg reflector (DBR) approach is presented using ZnMgSSe (instead of ZnSSe) layers as
the high-reflective-index material and MgS/ZnCdSe superlattices as the low index material. This
allows for the realisation of the stopband in the blue/violet spectral region from 400 to 460 nm.
Furthermore, this enables the employment of binary ZnSe QWs instead of ZnCdSSe QWs what leads
to a reduced inhomogenous broadening of the excitonic photoluminescence peak at elevated
temperatures.
When the Mg content of the quaternary ZnMgSSe layers is higher than 20%, a main challenge is to
achieve smooth DBR interfaces. Furthermore, the requirement of lattice matching to the GaAs
substrate needs precise control of deposition parameters during the DBR growth run. High-resolution
X-ray diffraction (HRXRD) measurements are performed for calibration of the composition. In order
to determine the exact quarter-wave thickness of each layer, the use of in-situ reflectometry turned out
to be crucial.
A 21 pair DBR with a stopband centered at 430 nm reaches a reflectivity exceeding 99%, while the
stopband width is about 40 nm. In Figures 1(a) and (b) reflectivity measurements and a scheme of the
new DBR concept are presented, respectively.
Vertical resonators formed by two DBR mirrors and a cavity containing binary ZnSe quantum wells
have been realized. In Figure 2(a) a vertical resonator with 18 bottom DBR pairs, a 1λ-cavity with
three ZnSe QWs and 15 top DBR pairs is presented. Micropillars have been etched out of the planar
structure using focused-ion-beam (FIB) milling (Figure 2(b)). The three-dimensional optical
confinement in the pillar microcavity results in the appearance of discrete modes which are observed
in the photoluminescence (PL) spectra (Figure 3) at the low-energy side of the QW emission. The
fundamental mode centered at 2.766 eV can be clearly identified for a pillar diameter of 1.46 μm. A
quality factor of 1600 can be determined from the spectral width of the fundamental mode.
In addition first results concerning the observation of strong coupling in these structures will be
discussed.
[1] C. Kruse et al., pss (b) 241, 731 (2004)
[2] P. Kelkar et al., Phys. Rev B 42 , R5491 (1995)
[3] J. Kasprzak et al., Nature 443, 409-414 (2005)
__________________________
* Contact: klembt@ifp.uni-bremen.de

Tu2.5
(a)
(b)
Fig 1: New II-VI DBR approach with ZnMgSSe as high index material, (a) reflectivity and refractive index comparison of
II-VI DBRs, (b) detailed structure of quaternary DBR.
(a)
(b)
Fig 2: Vertical resonator with a 18 pair bottom DBR, a 1λ-cavity with 3 ZnSe QWs and a 15 pair top DBR,
(a) schematic drawing, (b) micropillar with diameter of 1.46 μm.
Fig 3: PL spectrum of micropillar at T = 4 K. The spectral width of the fundamental mode shows Q = E/ΔE =1600.

Tuesday Session Tu3
Wide bandgap

Tu3.1
Recent device applications of non-polar cubic group III-Nitrides
D.J. As *
Universität Paderborn, Department Physik, Warburger Str. 100, 33098 Paderborn, Germany,
Group III-nitride-based optoelectronic and electronic devices, which are commercially
available at the market, are grown along the polar c direction, which suffer from the existence
of strong “built-in” piezoelectric and spontaneous polarization. This inherent polarization
limits the performance of optoelectronic devices containing quantum well or quantum dot
active regions. To get rid of this problem much attention has been focused on the growth of
non- or semi-polar (Al,Ga,In)N. However, a direct way to eliminate polarization effects is the
growth of cubic (100) oriented III-nitride layers. With cubic epilayers a direct transfer of the
existing GaAs technology to cubic III-Nitrides will be possible and the fabrication of diverse
optoelectronic devices will be facilitated. However, since cubic GaN is metastable and no
cubic GaN bulk material exists in nature, heteroepitaxy with all its drawbacks due to lattice
mismatch is necessary to grow this material. Due to the low lattice mismatch to cubic GaN the
substrate of choice for the growth of cubic III-nitrides is 3C-SiC.
In this talk the latest achievements in the molecular beam epitaxy of phase-pure cubic
GaN, AlN and their alloys grown on 3C-SiC substrates is reviewed [1, 2]. The structural and
optical properties of cubic nitrides, cubic AlGaN/GaN heterostructures and cubic GaN
quantum dots will be shown [3]. The absence of polarization fields in cubic nitrides is
demonstrated [4]. We summarize results of cubic AlGaN/GaN (c-AlGaN/GaN) heterojunction
field-effect transistors (HFETs) and introduce their output and transfer characteristics
with both normally-on and normally-off behaviour [5]. In highly n-doped bulk zincblende
GaN very long electron spin relaxation times exceeding 500 ps are observed up to roomtemperature,
which may be very interesting for future spintronic applications [6]. Furthermore
intersubband (ISB) absorption from superlattices of GaN/AlGaN with various quantum well
thicknesses shows absorption from 1.4 µm to 63 µm, covering the mid to far infrared range
and enters the THz regime [7]. This range also includes the technologically important
1.55 µm range for telecommunication, for which we also report first ISB inter-subband photo
responce measurements [8].
[1] D.J. As, Microelectronics Journal 40, 204 (2009)
[2] D.J. As, Proc. of SPIE Vol. 7608, 76080G (2010)
[3] T. Schupp, B. Neuschl, M. Feneberg, K. Thonke, K. Lischka, and D.J. As, Journal of Crystal Growth 312,
3235 (2010)
[4] J. Schörmann, S. Potthast, D.J. As, K. Lischka, Appl. Phys. Lett. 89, 131910 (2006)
[5] E. Tschumak, R. Granzer, J.K.N. Lindner, F. Schwierz, K. Lischka, H. Nagasawa, M. Abe, and D.J. As,
Appl. Phys. Lett. 96, 253501 (2010)
[6] J.H. Buß, J. Rudolf, T. Schupp, D.J. As, K. Lischka, and D. Hägele, Appl. Phys. Lett. 97, 062101 (2010)
[7] H. Machhadani, M. Tchernycheva, L. Rigutti, S. Saki, R. Colombelli, C. Mietze, D.J. As and F.H. Julien,
Phys. Rev. B (2010) (submitted)
[8] E.A. DeCuir, Jr., M.O. Manasreh, E. Tschumak, J. Schörmann, D.J. As, and K. Lischka, Appl. Phys. Lett. 92,
201910 (2008)
__________________________
* Contact: d.as@uni-paderborn.de

Tu3.2
Polar and nonpolar (Zn,Mg)O/ZnO heterostructures :
the benefits of homoepitaxy
J.-M. Chauveau 1,2 , M. Teisseire 1 , C. Morhain 1 , H. Chauveau 1 , C. Deparis 1 ,
B. Vinter 1,2
1
CNRS-CRHEA, Av. Bernard Grégory, F- 06560 Valbonne Sophia Antipolis, France
2 University of Nice Sophia Antipolis, Parc Valrose, F-06102 Nice Cedex 2, France
ZnO-based quantum wells have attracted much attention in the last few years due to their
opportunity of combining band gap engineering, with large excitonic binding energies. Indeed
theoretical works suggest that the 60meV-binding energy of excitons in ZnO could be further doubled
in quantum wells (QWs). So far studies on ZnO have mainly focused on films grown in (0001)
orientation. In this configuration, the wurtzite ZnO layers exhibit built-in electric fields (both piezo
and spontaneous components) along the c-axis, i.e. the growth direction, affecting the electronic
properties. Non-polar surfaces are therefore of a particular interest since the c-axis of the layer lies in
the growth plane in this case. As a result it is expected that (Zn,Mg)O/ZnO QWs structures can be
grown without any screening of the exciton binding energies.
A series of polar and nonpolar QWs with different widths and different Mg contents was
successfully grown by molecular beam epitaxy (MBE) on sapphire and ZnO substrates. A built-in
electric field of ~1MV/cm was observed in polar heterostructures [1] while nonpolar QWs exhibit the
absence of Quantum Confined Stark Effect [2] (Figure 1).
Wide band gap nonpolar QWs (nitrides or oxides) grown on sapphire usually exhibit a large
density of stacking faults, which strongly reduce the emission efficiency [3]. ZnO bulk substrates are
commercially available for polar and nonpolar orientations. Unfortunately, the as-received substrates
exhibited high densities of scratches at the surface. We show that a dedicated surface preparation is
needed before MBE growth. Indeed prior to the growth, the ZnO substrates were in-situ en ex-situ
prepared and streaky RHEED patterns were achieved. The atomically flat surfaces were obtained after
the annealing procedure (Figure 2a).
In this presentation we shall compare hetero- and homo-epitaxial QWs. We show a drastic
improvement of the structural properties when the QWs are grown on ZnO substrates: no residual
strain in QWs for thin (Zn,Mg)O barriers, surface roughness and X-Ray full width at half maximum
(FWHM) reduced by a factor of ten [4]. The high resolution transmission electron microscopy images
exhibit smooth interfaces between the (Zn,Mg)O barriers and the QW, without extended defects
(Figure 2b).
The drastic improvement of the structural properties implies a strong enhancement of the
photoluminescence (PL) properties. The PL FWHM can be as low as 3.5meV at 8K for QWs larger
than 4nm while the intensity is about 50 times larger than that of QWs grown on sapphire (Figure
3a).… For this series of QWs, the free exciton with the lowest energy was measured by
photoluminescence excitation (PLE). Its position was calculated by taking into account the variation of
the exciton binding energy with the QWs width [5]. Note the very good agreement with the PLE data
(figure 3b). It is also remarkable that the RT PL emission of the nonpolar homoepitaxial a-plane QW is
still one order of magnitude more intense than that of the nonpolar heteroepitaxial QW taken at low
temperature. Finally we shall compared the different nonpolar orientations (m-plane or a-plane) and
we show that the PL intensity of an m-plane QW is constant as a function of the temperature up to RT,
indicating the thermal equilibrium of excitons.
The comparison between hetero- and homo- epitaxial QW heterostructures convincingly
demonstrates the interest of homoepitaxial QWs for bright UV emission applications.

Tu3.2
Fig 1: PL emission dependence of ZnO/(Zn,Mg)O QWs recorded at 10K for various widths of the quantum wells
(Lw). Two orientations are compared: QWs grown along the c direction (open triangles, Ref. [1]) and QWs grown in
the a-direction (square, [2]). The dotted line shows the exitonic gap of ZnO.
600nm
(a)
Fig 2: (a) Surface of a ex-situ prepared ZnO substrate, (b) Cross sectional Cs corrected high resolution TEM image
11 00 zone axis. Both quantum well and barriers are
indicated.
taken from the 3.5 nm (22 monolayers) QW sample along the [ ]
(b)
PL Intensity
6000
QW on ZnO
4000
2000
QW on sapphire
x50
XA position (eV)
3.57
3.50
3.43
calculation
PLE
12
10
8
6
4
2
FWHM (meV)
3.35 3.40 3.45 3.50
Energy (eV)
3.36
0
1 2 3 4 5 6 7 8
QW Thickness (nm)
(a)
Fig 3: (a) Comparison between PL spectra from homoepitaxial and heteroepitaxial QWs. (b) Ground state energies
(PLE) from the series of QWs (open squares) taken at 2K. Calculated transitions corrected by the exciton binding
energy (open circles). The error bars are deduced from the FWHM measured in PL on the energy and +/- one
monolayer on the QW thickness. The FWHM is indicted by full stars.
__________________________
* Contact: jmc@crhea.cnrs.fr
(b)
1 C. Morhain, T. Bretagnon, P. Lefebvre, X. Tang, P. Valvin, T. Guillet, B. Gil, T. Taliercio,
M. Teisseire-Doninelli, B. Vinter, and C. Deparis, Phys. Rev. B 72 (24), 241305 (2005).
2 J. M. Chauveau, D. A. Buell, M. Laugt, P. Vennegues, M. Teisseire-Doninelli, S. Berard-
Bergery, C. Deparis, B. Lo, B. Vinter, and C. Morhain, J. Cryst. Growth 301, 366 (2007).
3 P. Vennegues, J. M. Chauveau, M. Korytov, C. Deparis, J. Zuniga-Perez, and C. Morhain,
J. Appl. Phys. 103 (8), 083525 (2008).
4 J.-M. Chauveau, M. Teisseire, H. Kim-Chauveau, C. Deparis, C. Morhain, and B. Vinter,
Appl. Phys. Lett. 97, 081903 (2010).
5 J. M. Chauveau, J. Vives, J. Zuniga-Perez, M. Laugt, M. Teisseire, C. Deparis, C.
Morhain, and B. Vinter, Appl. Phys. Lett. 93 (23), 231911 (2008).

Tu3.3
A bi-layer oxide buffer approach for the integration of single
crystalline GaN on Si (111) platform
L. Tarnawska 1,* , P. Zaumseil 1 , M. Kittler 2 , P. Storck 3 , R. Paszkiewicz 4 , and T. Schroeder 1
1 IHP, Im Technologiepark 25, 15236 Frankfurt (Oder), Germany
2 IHP/ BTU Joint Lab, Konrad-Wachsmann-Allee 1, 03013 Cottbus, Germany
3 SILTRONIC AG, Hanns-Seidel-Platz 4, 81737 München, Germany
4 Wroclaw University of Technology, Janiszewskiego 11/17, 50-372 Wroclaw, Poland
Integration of GaN virtual substrates on Si wafers is intensively pursed for high power-high frequency
electronics as well as optoelectronics applications. However, the growth of GaN layers on Si substrates
presents several difficulties related to the very high reactivity of the Si surface with nitrogen, the large
lattice mismatch (-17%), and the large difference in thermal expansion coefficient (33%). As a
consequence GaN epitaxial layers grown on Si show a large number of threading dislocations density
severely deteriorating the overall quality of the GaN films. To overcome these problems by
accommodating the misfit and avoiding interfacial reactions, different semiconducting (e.g AlN, GaAs,
ZnO) and insulating buffer layers (e.g. γ-Al 2 O 3 ) were used in the past. In this work, we present a novel bilayer
buffer approach for the integration of GaN on Si(111) via Sc 2 O 3 /Y 2 O 3 intermediate layer system.
The samples were grown in a multichamber molecular beam epitaxy (MBE) system. Boron-doped 4-inch
Si(111) wafers were used as substrates. Substrate temperature during growth of Y 2 O 3 and Sc 2 O 3 was 650
and 500°C, respectively. The growth process was in-situ monitored by RHEED. After deposition of the
oxide layers samples were transferred under UHV to another MBE chamber for plasma assisted GaN
overgrowth. Deposition conditions for GaN were optimized to achieve best quality with respect to surface
morphology, crystalline quality, optical and electrical properties [1]. To obtain complete information on
the quality of GaN/Sc 2 O 3 /Y 2 O 3 /Si(111) heterostructures, RHEED, XPS, XRR and XRD, SEM and TEM
measurements were complemented by Synchrotron Radiation Grazing Incidence (SR-GI) XRD.
Our studies demonstrate that the Sc 2 O 3 /Y 2 O 3 buffer approach on Si(111) provides a template of high
structural quality for GaN overgrowth [2]. Buffer oxides are of type-B stacking with respect to the
Si(111) substrate and are fully relaxed. The 7% lattice mismatch between the oxides is compensated by
edge dislocation formed by insertion of additional {111} planes in the Sc 2 O 3 thin film. In addition, Sc 2 O 3
and Y 2 O 3 are thermodynamically stable at least up to 900°C what is sufficient for GaN growth by MBE.
Figure 1 shows a low magnification TEM image (a) and specular θ-2θ XRD scan (b) for a ~ 30 nm thick
GaN layer. These measurements confirm that the growth of GaN(0001)/Sc 2 O 3 (111)/Y 2 O 3 (111)/Si(111)
heterostructure is achieved. Temperature-dependent XRD measurements were performed to determine the
coefficients of thermal expansion of the buffer oxide and GaN film (Fig. 1(c)). The in-plane orientation
was studied by SR-GI XRD and it was found to be GaN[10-10]||Sc 2 O 3 [2-1-1]||Y 2 O 3 [2-1-1]||Si[-211].
Since the GaN film quality is dependent on the quality of the initial nucleation layer, in-situ growth
studies by RHEED and XPS are under way to reveal the nature of the atomic bonding between Sc 2 O 3 and
GaN. To optimize GaN growth conditions, 600 nm-thick layers were grown and characterized. The main
defects found in thicker GaN layers are threading dislocations (TD), with density in the order of 10 10 cm -2 ,
and stacking faults, resulting in microscopic cubic grains within the hexagonal matrix.
Cathodoluminescence show a near-band-edge emission at 358 nm and a bread yellow luminescence band
at 560 nm (Fig. 1(d)).
[1] H. Markoç, J. Material Science, 12, 677 (2001)
[2] L. Tarnawska, A. Giussani, P. Zaumseil, M.A. Schubert, R. Paszkiewicz, O. Brandt, P. Storck, T. Schroeder, J. Appl.
Phys.,108, 1 (2010)
__________________________
*Contact: tarnawska@ihp-microelectronics.com

Tu3.3
2.61 average CTE = 4.1x10-6 1/K
c GaN
2.60
after cooling
c [A]
2.59
(a)
2.58
0 200 400 600 800
Temeprature [°C]
(c)
Si 222
7x10 4
near BE
358 nm
intensity [cps]
10 4
10 3
Y Sc GaN 0004
2
O 3
444
2
O 3
444
10 2
10 1
55 60 65 70 75 80
2θ [degree]
CL intensity [a.u.]
(b)
6x10 4
5x10 4
YL
560 nm
300 400 500 600
λ [nm]
(d)
Fig.1: (a) TEM cross -section of the GaN(0001) / Sc 2 O 3 (111) / Y 2 O 3 (111) / Si(111) heterostructure, viewed along a
azimuth; (b) specular θ-2θ XRD measurement; (c) extraction of the coefficient of thermal expansion for GaN thin layer; (d)
Cathodoluminescence for 600 nm-thick GaN film

Tu3.4
GaN/AlGaN superlattices grown by PAMBE for intersubband
applications in the infrared spectral range
Y. Kotsar 1, * , A. Das 1 , E. Bellet-Amalric 1 , E. Sarigiannidou 3 H. Machhadani 2 , S.
Sakr 2 , M. Tchernycheva 2 , F. H. Julien 2 and E. Monroy 1
1
CEA-CNRS group « Nanophysique et semiconducteurs », INAC/SP2M/NPSC,
CEA-Grenoble, 17 rue des Martyrs, 38054 Grenoble cedex 9, France
2
Institut d’Electronique Fondamentale, UMR 8622 CNRS, Université Paris-Sud, 91405 Orsay cedex, France
3
LMGP, Grenoble INP, 3 Parvis Louis Néel, BP 257, 38016 Grenoble cedex 1, France
The intersubband (ISB) technology, based on electronic transitions between the confined levels in
the conduction or valence band of semiconductor nanostructures, offers the possibility to tune the
operation wavelength by changing the design of the active region. GaN/Al(Ga)N heterostructures are
of a great interest for ISB device applications due to the large conduction band offset (1.75 eV
between GaN and AlN), which allows building devices that cover the near- and mid-infrared spectral
regions. In addition, the short ISB relaxation time (≈150 fs) via LO-phonon scattering process finds
application in ultra-high-speed quantum cascade photodetectors and electro-optical/ all-optical
modulators. Moreover, the large LO-phonon energy (92 eV) of GaN should enable room-temperature
operation of quantum cascade lasers in the far-infrared spectral range.
The active region of ISB devices consists basically of GaN/Al(Ga)N QW or coupled-QW
superlattices (SLs). In such structures, strain management is important due to the large lattice
mismatch between GaN and AlN (2.4%). The strain relaxation affects the quality of the epitaxial
layers by introducing structural defects and modifying the SL band profile via piezoelectric effects,
which in turn can influence the confined energy levels. Therefore, it is essential to account for growth
aspects, stress generation and strain relaxation mechanisms when designing the active medium.
We have investigated the relaxation mechanism of 40-period 7 nm/4 nm GaN/Al x Ga 1-x N SLs
(x = 0.1, 0.3, 0.5), designed for mid-infrared ISB absorption in the 5-10 µm spectral range, and
deposited by plasma-assisted molecular-beam epitaxy (PAMBE). To have a better understanding of strain
relaxation, we consider samples grown either on GaN or on an AlGaN buffer layer with the same Al
mole fraction than the SL barriers. In situ and ex situ studies of the grown structures were performed
(Fig. 1 and 2). In the case of growth on GaN and for x = 0.1, 0.3 we observe a pseudomorphic growth
favored by the surface energy minimization effect under Ga excess growth conditions [Fig. 1(a) and
Fig. 2(a)]. In contrast, higher Al mole fraction or growth on AlGaN results in a gradual plastic
relaxation during the first 10-20 periods of the SL, with the corresponding increase of the density of
threading dislocations with a Burgers vector b = ⅓‹11-20› [Fig.2 (b)]. In addition to the average
relaxation trend of the SL, RHEED measurements indicate a periodic relaxation. In the case of
GaN/AlN, we have demonstrated that this relaxation is associated to a plastic deformation with
appearance of dislocations associated to stacking fault loops. In the case of GaN/Al x Ga 1-x N (x = 0.1,
0.3, 0.5), we demonstrate that the periodic fluctuation of in-plane lattice parameter can be attributed to
an elastic phenomena due to the stress induced by the metal excess on the growing surface (Fig. 3).
The potential of these structures for infrared optoelectronics is discussed. GaN/AlGaN SLs
displaying ISB absorption in the 1.3-10 µm near- and mid-infrared spectral range are presented, and
spectroscopic results are compared with theoretical calculations of the electronic structure. For the farinfrared
spectral range, beyond the Reststrahlen band of GaN (13-20 µm), we have investigated
heterostructures with a flattened band profile using GaN/Al 0.05 Ga 0.95 N/Al 0.1 Ga 0.9 N step quantum wells.
In these structures we observed low-temperature ISB absorption at 71.4 µm and 142.9 µm which sets a
new wavelength record for this material system (Fig. 4).

Tu3.4
3.200
Relaxed GaN
3.200
Relaxed GaN
Lattice parameter (Å)
3.180
3.160
3.140
3.120
Relaxed Al 0.1 Ga 0.9 N
Relaxed AlN
(a)
Lattice parameter (Å)
3.180
3.160
3.140
3.120
Relaxed Al 0.5 Ga 0.5 N
Relaxed AlN
(c)
3.100
0 5 10 15 20 25 30 35 40
SL Period
3.100
0 5 10 15 20 25 30 35 40
SL Period
Fig 1: Relaxation path traced through evolution of the in-plane lattice parameter along 40-period GaN/Al x Ga 1-x N (7
nm / 4 nm) SLs grown either on Al x Ga 1-x N or on GaN for (a) x = 0.1, (b) x = 0.5. The lattice parameters of relaxed
AlN, Al x Ga 1-x N and GaN are indicated by dashed lines.
(a)
(b)
Fig 2: TEM images of GaN/Al 0.1 Ga 0.9 N SLs taken near the ‹1-100› zone axis (a, b) with diffraction vector g = (11-20)
of structure strained on (a) GaN and (b) on Al 0.1 Ga 0.9 N buffer layers.
∆a/a (%)
0.2
0.1
0.0
-0.1
Ga cell open
Ga cell close
x = 0
x = 0.1
x = 0.3
x = 0.5
x = 1
-0.2
-0.3
0 20 40 60 80 100 120 140
Time (s)
Fig 3: Evolution of in-plane lattice parameter induced by
depositing Ga on top of GaN, AlN and Al x Ga 1-x N.
Fig 4: Transmission spectra for TM- (square) and TE-
(circle) polarized light at 4.7 K with a peak absorption at
2.1 THz.
__________________________
* Contact: yulia.kotsar@cea.fr

Tu3.5
Growth and Characterization of ZnO/ZnMgO Quantum Wells
Bernhard Laumer 1, 2,* , Fabian Schuster 1 , Thomas Wassner 1 , Martin Stutzmann 1 , and
Martin Eickhoff 2
1 Walter Schottky Institut, Technische Universität München, Am Coulombwall 3, 85748 Garching, Germany
2 I. Physikalisches Institut, Justus-Liebig-Universität, Heinrich-Buff-Ring 16, 35392 Giessen, Germany
Hetero- and homoepitaxial ZnO/Zn 1-x Mg x O single quantum wells (SQWs) with different quantum well
widths d W and barrier compositions up to x = 0.27 were grown by plasma-assisted molecular beam
epitaxy on a-plane and c-plane sapphire as well as homoepitaxially on (000-1)- and (11-20)-oriented
ZnO substrates. For heteroepitaxy on c-plane sapphire, an MgO/ZnO double buffer has been
introduced to overcome the large lattice mismatch [1]. In order to inhibit phase separation into
wurtzite Zn(Mg)O and cubic Mg(Zn)O for high Mg concentrations, two different approaches were
pursued: First, graded Zn 1-x Mg x O barriers with a maximum Mg content x = 0.24 were realized by a
step-wise increase of the Mg concentration x in consecutive, thin Zn 1-x Mg x O layers. In the second
approach, a low substrate temperature of T S = 270 °C and metal-rich conditions were adopted to
achieve Zn 1-x Mg x O barriers with an Mg content of x = 0.27.
Atomic force microscopy reveals smooth surfaces with an rms roughness between 0.3 and 0.4 nm for
both methods, although atomic steps are only observable for growth temperatures above 400°C. Time
of flight secondary ion mass spectroscopy was employed to investigate the abruptness of the heterointerfaces
and the occurrence of intermixing for different growth conditions and substrates. Reciprocal
space maps of graded Zn 1-x Mg x O barriers grown on c-plane sapphire reveal a partial relaxation of the
individual Zn 1-x Mg x O layers, while fully relaxed growth is observed on a-plane sapphire.
The optical properties were investigated by photoluminescence spectroscopy. At 4.2 K a maximum
blue shift of 280 meV of the SQW emission with respect to bulk ZnO is observed for d W = 0.8 nm and x =
0.27. Temperature-dependent measurements reveal the highest temperature stability of the SQW
emission for d W = 1.2 nm with an activation energy of (95 ± 5) meV for thermal decay. For polar
samples with x ≥ 0.16 and sufficiently broad wells, the SQW emission drops below that of bulk ZnO,
unambiguously evidencing the influence of the quantum confined Stark effect. Furthermore, the SQW
emission is quickly quenched for temperatures above 120 K and exhibits pronounced LO phonon
replicas. Increasing the excitation power results in partial screening of polarization charges, which in
turn causes a blue shift of the SQW emission. Numerical simulations reveal an internal electric field of
approximately 700 kV/cm for x = 0.27 (Fig. 1). In contrast, in non-polar SQWs grown
homoepitaxially on (11-20)-oriented ZnO substrates no internal electric fields are present and therefore
an enhanced thermal stability of the SQW emission as compared to polar samples is established.
[1] A. Bakin et al., J. Crystal Growth, 287, 7 (2006)
__________________________
* Contact: bernhard.laumer@wsi.tum.de

Tu3.5
Intensity [arb. units]
QW emission
bulk ZnO
barrier
0.8nm
1.2nm
1.8nm
2.3nm
3.5nm
7.0nm
Transition Energy [eV]
3.7
3.6
3.5
3.4
3.3
3.2
3.1
3.0
E int
= 0 kV/cm
E int
= 700 kV/cm
Experimental data
bulk ZnO
9.0nm
2.9
(a)
3.0 3.2 3.4 3.6 3.8 4.0 0 2 4 6 8 10
Energy [eV]
(b)
Well Width [nm]
Fig. 1: (a) Low-temperature PL spectra of ZnO/Zn 0.73 Mg 0.27 O SQWs with various well widths. The SQW emission is marked
by arrows. (b) Comparison of experimental transition energies with simulations without an internal field E int and for E int =
700 kV/cm.

Tu3.6
InGaN laser diodes operating at 450-455 nm grown by RF-Plasma MBE
M Siekacz 1,2,* , H. Turski 1 , M. Sawicka 1,2 , G. Cywiński 1 , J. Smalc-Koziorowska 1,2 ,
P. Wiśniewski 1,2 , P. Perlin 1,2 , I. Grzegory 1 and C. Skierbiszewski 1,2
1
Institute of High Pressure Physics, Polish Academy of Sciences, 01-142 Warszawa, Poland
2
TopGaN Ltd, ul Sokolowska 29/37, 01-142 Warszawa, Poland
Growth of high indium content structures for blue and green lasers diodes (LDs) is recently
one of the biggest challenges for epitaxy of nitrides. Due to the extremely large electrical polarization
contrast between high In-content quantum wells and gallium nitride (related to the compressive strain)
and also with the difficulty of growing homogenous InGaN having over 20% of In, the fabrication of
green light emitting structures is very demanding. Very recently, green LDs at 500-530 nm have been
demonstrated in nitride-based structures grown by Metal Organic Vapour Phase Epitaxy (MOVPE)
either on polar, semipolar and nonpolar substrate orientations. On the other hand, progress in
understanding the new growth mechanism for nitrides in Plasma Assisted Molecular Beam Epitaxy
(PAMBE) has led to the demonstration of blue-violet laser diodes, which in turn has renewed interest
in this technology among the MBE community. For PAMBE, the growth mechanism is entirely
different from that in MOVPE, and allows for the growth of device-quality nitride structures at
temperatures lower by 200-300°C versus those used in MOVPE. Therefore it is highly interesting
whether this technology can be useful for high In content structures required for green emitters. We
have shown already last year that by increasing the nitrogen flux in PAMBE we were able to
demonstrate optically pumped lasing from Single Quantum Well (SQW) InGaN laser structures in the
range 470 - 501 nm.
In this work we demonstrate first blue laser diodes grown by MBE which operate at the region
of 450 - 455 nm. Laser diodes were grown on c-plane GaN HVPE substrates, with dislocations (TDs)
density 10 6 -10 7 cm -2 . The laser structures consist of staggered SQW InGaN. In staggered SQW, the
indium content is increased in two steps, therefore piezoelectric field in the active part of SQW is
smaller than in standard SQW of the same In content. The devices were processed as ridge-waveguide,
oxide-isolated lasers. The mesa structure 10 µm x 1400 µm was etched out in the wafer down to the
depth of 0.3 µm. We observe only 10 nm blue shift between lasing emission and a spontaneous one -
see Figure 1. The key element to achieve lasing at 455 nm was to reduce the light losses and to keep
optical modes inside waveguide. We will discuss the influence of InGaN waveguides and heavily Si
doped GaN plasmonic claddings to LDs performance (see Fig. 2 for LDs structure details). The
Atomic Force Microscopy, Transmission Electron Microscopy and selective defect etching were used
to determine structural quality of LDs. The number of TDs after the growth of LDs was the same as in
substrate. We will present also results of theoretical calculations of light losses in LDs as a function of
the (i) claddings thickness and Al content, (ii) thickness and doping of GaN:Si plasmonic cladding (iii)
the composition of In in waveguide.
This work show that present challenges for long wavelength nitride laser diodes are very
similar for both technologies: PAMBE and MOVPE. At the moment for PAMBE the most important
task is to reduce the optical losses for long wavelength devices. Due to the smaller contrast of GaN
and AlGaN refractive indexes and much lower optical gain for structures operating at 450-500 nm, the
construction of LDs is more demanding in comparison to blue LDs. We show that PAMBE is an
alternative technology for MOVPE for blue-green optoelectronic devices.
Acknowledgements: This work was supported partially by the Polish Ministry of Science and Higher Education Grant No IT
13426 and the European Union within European Regional Development Fund, through grant Innovative Economy
(POIG.01.01.02-00-008/08) and SINOPLE project.
*
Corresponding author: Marcin Siekacz, Institute of High Pressure Physics, Polish Academy of Sciences, Sokołowska 29/37,
01-142 Warsaw, Poland, Phone: +48 22 6325010, Fax: +48 22 6324218, email: msiekacz@unipress.waw.pl

Tu3.6
Optical Power (mW)
40
30
20
LD PAMBE
pulse operation
10
5
(a)
0
0 2 4 6 8 10 12 14 16 0
Current Density (kA/cm 2 )
30
25
20
15
10
Voltage (V)
Light intensity (a.u.)
3
2
1
(b)
λ=455.4 nm
J TH
= 11 kA/cm 2
0
420 440 460 480 500 520
λ (nm)
Figure 1. (a) L-I-V characteristic of PAMBE laser diode, (b) the spontaneous and stimulated emission of this diode.
Figure 2. Laser diode structure details

STANDARD EFFUSION CELLS
ORGANIC MATERIAL EFFUSION
CELLS
LOW TEMPERATURE EFFUSION
CELLS
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UHV E-BEAM EVAPORATORS
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EVAPORATORS
VALVED CRACKER SOURCES
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HYDROGEN ATOM BEAM SOURCES
OXYGEN ATOM BEAM SOURCES
SUBSTRATE MANIPULATORS
CUSTOMIZED EQUIPMENT
AsBr 3
IN- SITU ETCHING SYSTEMS
MBE-SYSTEMS
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Wednesday Session We1
Group IV Materials

We1.1
Surfactant-modified epitaxy of germanium layers on silicon for
high-mobility channels
T. F. Wietler *
Institute of Electronic Materials and Devices, Leibniz Universität Hannover,
Schneiderberg 32, 30167 Hannover, Germany
The incorporation of germanium as a high-mobility channel material into silicon-based metal oxide
semiconductor field-effect transistor (MOSFET) technology is widely accepted as a solution to adhere
to Moore’s law under aggressive scaling beyond 32 nm. Germanium’s higher carrier mobility could
compensate for the performance degradation anticipated for silicon channel devices [1-5]. To achieve
this goal, the integration of epitaxial germanium layers is a must.
Heteroepitaxy of germanium on silicon is hindered by the 4.2% lattice mismatch that leads to threedimensional
islanding and heavy generation of film penetrating defects which deteriorate device
performance. One way to overcome this obstacle is the use of surfactants [6]. Surfactant-mediated
epitaxy (SME) has proven to enable the growth of smooth fully relaxed germanium layers with low
defect densities on Si(1 1 1) [7,8] and also on Si(0 0 1) substrates [9].
This paper will briefly outline the different approaches towards the integration of epitaxial germanium
into silicon based devices and then, focus on the recent progress in SME of germanium on silicon
substrates. In particular, insight will be given into the antimony controlled strain relaxation
mechanism, which is crucial for the structural properties of the resulting films. At growth temperatures
above 600°C, the presence of antimony at the growth front suppresses three-dimensional islanding and
instead leads to the development of a regularly facetted surface. On Si(1 1 1) the micro-facets are
oriented while on Si(0 0 1) their shape depends on antimony coverage, being for low
coverage and for full coverage. It has been shown, that the facet orientation and, in particular,
the direction of the troughs between these facets is crucial for the strain relaxation process during
further growth [10]. The troughs are locations with excess strain, and thus provide preferred nucleation
sites for misfit dislocations. Only troughs with directions in the surface result in regular
dislocation networks confined at the interface allowing for smooth fully relaxed germanium films with
low dislocation densities.
Given the excellent structural properties of relaxed germanium films grown by SME, the issue of
surfactant incorporation induced background doping has been studied. While germanium films grown
at temperatures around 500°C suffer from high background doping (~10 19 cm -3 ), elevated growth
temperatures enhance the surface segregation of antimony. This enables low background doping and
consequently, high electron and hole mobilities necessary for device purpose.
In view of germanium channel MOSFETs, the challenging quest for a perfect gate dielectric on
germanium will also be addressed with special emphasis on epitaxial high-k materials. As an outlook,
possible applications of epitaxial germanium films on silicon substrates beyond high-mobility
channels will be broached.
[1] D. A. Antoniadis, 2002 Symp. VLSI Technol. Digest of Technical Papers., 2 (2002).
[2] D.A. Antoniadis, I. Aberg, C. Ní Chléirigh, O.M. Nayfeh, A. Kakifirooz, and J.L. Hoyt., IBM J. Res. & Dev., 50 363
(2006).
[3] M. Lundstrom, Science, 299 210 (2003).
[4] R. Chau, B. Doyle, S. Datta, J. Kavalieros, and K. Zhang, Nature Materials, 6 810 (2007).
[5] H. Shang, M.M. Frank, E.P. Gusev, J.O. Chu, S.W. Bedell, K.W. Guarini, and M. Ieong, IBM J. Res. & Dev., 50 377
(2006).
[6] M. Copel, M.C. Reuter, E. Kaxiras, and R.M. Tromp, Phys. Rev. Lett., 63 632 (1989).
[7] M. Horn-von Hoegen, F.K. LeGoues, M. Copel, M.C. Reuter, and R.M. Tromp, Phys. Rev. Lett., 67 1130 (1991).
[8] T.F. Wietler, A. Ott, E. Bugiel, and K.R. Hofmann, Mat. Sci. Semicond. Proc., 8 73 (2005).
[9] T.F. Wietler, E. Bugiel, and K.R. Hofmann, Appl. Phys. Lett., 87 182102 (2005).
[10] T. F. Wietler, E. P. Rugeramigabo, E. Bugiel, and K. R. Hofmann. AIP Conf. Proc. 1199 (2009) 15–16.
__________________________
* Contact: wietler@mbe.uni-hannover.de

We1.2
Epitaxial growth of SrTiO 3 on Si : strain relaxation and formation
of tetragonal domains
G. Saint-Girons 1* , G. Niu 1 ; J. Penuelas 1 , L. Largeau 2 , B. Vilquin 1 , J.L. Maurice 3 C.
Botella 1 and G. Hollinger 1 .
1 Université de Lyon, Institut des Nanotechnologies de Lyon (UMR5270/CNRS), Ecole Centrale de Lyon, 36
avenue Guy de Collongue, 69134 Ecully cedex, France
2 LPN-UPR20/CNRS, Route de Nozay, 91460 Marcoussis, France.
3 LPICM CNRS/Ecole Polytechnique, 91128 Palaiseau cedex, France
The epitaxial growth of SrTiO 3 (STO) on Si has been quite intensively studied in the last years
[1,2,3,4,5,6], because it opens the unique opportunity of integrating functional oxides or even III-V
materials [7,8,9]on Si. In particular, the way how Si surface can be passivated in a molecular beam
epitaxy (MBE) environment for further STO growth and the impact of the passivation process on
interface abruptness is now clearly described. However, STO structural quality is still far from
perfection, and much remains to be done to fabricate reliable STO/Si templates and to understand the
STO growth mechanisms. As an example, a recent study reports the observation of room-temperature
ferroelectricity of thin STO/Si layers associated with a significant tetragonality of the STO lattice,
without clearly describing the origin of this behavior [10].
In the present contribution, we will show that in specific growth conditions, thin STO/Si layers are
two-phased. On the basis of transmission electron microscopy (TEM); X-Ray diffraction (XRD) and
reflection high energy electron diffraction (RHEED) experiments, we will show that
-Strained cubic STO and anomalously tetragonal STO coexists in thin STO/Si layers
-The strained cubic STO phase remains fully coherent to Si up to at least 24 ML
-The tetragonal STO phase is not coherent to Si and progressively transits to cubic STO above 24 ML.
The first few STO monolayers (ML) have to be grown on Si at quite moderated temperature to avoid
interface reactions. We will show that this leads to an initially partly amorphous STO growth at the
origin of the formation of the two phases.
On the basis of these results, we will discuss the origin of ferroelectricity in ultrathin STO/Si layers,
and we will propose a growth process that allows fabricating single phased STO layers have substrate
like structural quality.
Fig.1 : TEM characterizations of a thin STO/Si layer providing evidence for the presence of two STO phases : a well
ordered cubic STO phase (zone 1 in (a)), and a more disordered tetragonal phase (zone 2 in (a))
__________________________
* Contact: Guillaume.saint-girons@ec-lyon.fr

We1.2
Fig.2 : Evolution of the in and out-of-plane lattice parameters of both STO phases as deduced from X-Ray diffraction
experiments. Cubic STO is fully strained on Si up to 24 ML, and its critical thickness is between 24 and 42 ML.
Tetragonal STO is not coherent to Si and transits to cubic STO when the deposited thickness exceeds 24 ML.
[1] R.A. McKee, F.J. Walker and M.F. Chisholm, Phys. Rev. Lett. 81, 3014, (1998).
[2] R.A. McKee, F.J. Walker and M.F. Chisholm, Science 293, 468, (2001).
[3] X. F. Wang, J. Wang, Q. Li, M. S. Moreno, X. Y. Zhou, J. Y. Dai, Y. Wang and D. Tang, J. Phys. D: Appl. Phys. 42,
085409 (2009)
[4] H. Li, X. Hu, Y. Wei, Z. Yu, X. Zhang, R. Droopad, A. A. Demkov, J. Edwards, Jr., K. Moore, and W. Ooms, J. Kulik and
P. Fejes, J. Appl. Phys. 93, 4521 (2003)
[5] G. Delhaye, C. Merckling, M. El-Kazzi, G. Saint-Girons, M. Gendry, Y. Robach, and G. Hollinger, L. Largeau and G.
Patriarche, J. Appl. Phys. 100, 124109 (2006)
[6] S. B. Mi, C. L. Jia, V. Vaithyanathan, L. Houben, J. Schubert, D. G. Schlom, and K. Urban, Appl. Phys. Lett. 93, 101913
(2008)
[7] G. Saint-Girons, P. Regreny, L. Largeau, G. Patriarche and G. Hollinger, Appl. Phys. Lett. 91, 241912, (2007).
[8] G. Saint-Girons, C. Priester, P. Regreny, G. Patriarche, L. Largeau, V. Favre-Nicolin, G. Xu, Y. Robach, M. Gendry, and G.
Hollinger, Appl. Phys. Lett. 92, 241907 (2008)
[9] G. Saint-Girons, J. Cheng, P. Regreny, L. Largeau, G. Patriarche and G. Hollinger, Phys. Rev. B 80, 155308 (2009)
[10] M. P. Warusawithana, C. Cen, C. R. Sleasman, J. C. Woicik, Y. Li, L. F. Kourkoutis, J. A. Klug, H. Li, P. Ryan, L. P.
Wang, M. Redzyk, D. A. Muller, L. Q. Chen, J. Levy and D. G. Schlom, Science, 324, 367 (2009)
__________________________
* Contact: Guillaume.saint-girons@ec-lyon.fr

We1.3
Morphology and luminescence properties of Sb mediated Ge/Si
quantum dots
A.A.Tonkikh 1,2,* , N.D.Zakharov 1 , V.G.Talalaev 3 , A.V.Novikov 2 , K.Kudryavtsev 2 ,
B.Fuhrmann 4 , H.S.Leipner 4 , P.Werner 1
1
Max-Planck Institute of Microstructure Physics, Weinberg 2, Halle, Germany
2 Institute for Physics of Microstructures RAS, Nizhniy Novgorod, Russia
3 ZIK SiLi-nano, Martin Luther University, Halle, Germany
4 Interdisciplinary Center of Materials Science, Martin Luther University, Halle, Germany.
There is currently a considerable interest in development of an efficient silicon based light
emitting source aimed to fill the active device niche for chip to chip communications [1]. The hybrid
III-V on Si technology dominates, but other approaches are intensively investigated. Among of them
the approach of small Ge inclusions grown epitaxially inside a Si p-n junction in the form of selfassembled
islands referred to as quantum dots (QDs) is considered. In spite of the fact that the bulk Ge
itself has an indirect band gap the energy band diagram of the Ge QD is modified due to the elastic
strain and the quantum confinement effect. In particular, high density of Ge QDs leads to the higher
luminescence probability of ∆ xy – HH transitions [2]. Here ∆ xy are electronic states localized in the
compressively strained Si between the QDs, HH is the heavy hole state in the QD.
Since the kinetics governs Ge QDs formation during MBE, the lower the growth temperature
the higher the QD density is expected. However the quality of the epi-film is strongly affected by the
growth temperature. In practice this means that the ability to control the QD array properties is limited.
To bypass this difficulty the surfactant assisted QD growth can be applied.
A thin Sb layer was shown to reduce the Ge ad atom migration ability leading to the formation
of smaller QDs [3]. We utilized and developed this technique further to obtain an array of the small Ge
QDs at elevated (>550 0 C) temperatures. The submonolayer Sb film was deposited before the Ge QDs
growth. In the case of uncapped Ge islands the morphology was investigated by the atomic force
microscopy. The Ge QDs embedded in the Si host were analyzed using the transmission electron
microscopy. It was found that the Ge islands separated by about 11 nm from each other, having sizes
of about 15 nm (basis) and 2.5 nm (height) were formed at the temperature of 600 0 C on the Si (100)
substrates. We have shown that the Sb mediated Ge QDs array capped with the Si layer exhibited room
temperature photoluminescence at a wavelength of about 1.5 µm [4].
The present contribution discusses the growth mechanism of the Sb mediated Ge QDs, their
luminescence properties and the size effects. Electroluminescence data on the Si p-n junction with the
embedded arrays of the small Ge QDs will be presented as well.
[1] ITRS: 2009 (http://www.itrs.net/Links/2009ITRS/Home2009.htm).
[2] M Brehm et. al. New J. Phys. 11, 063021 (2009).
[3] A. Portavoce, I. Berbezier, and A. Ronda, Phys.Rev.B 69, 155416 (2004).
[4] A.Tonkikh et. al. Phys.Stat.Sol. RRL 4, 224 (2010).
__________________________
* Contact: tonkikh@mpi-halle.mpg.de

We1.4
Growth of small-period Si/Ge quantum dot crystals by MBE
S. Borisova 1,* , C. Dais 2 , J.C. Gerharz 1 , G. Mussler 1 , and D. Grützmacher 1
1
Institute of Bio- and Nanosystems 1, Forschungszentrum Jülich, D-52425 Jülich, Germany
2 Laboratory for Micro- and Nanotechnology, Paul Scherrer Institut, CH-5232 Villigen-PSI, Switzerland
For decades, silicon has been the most ubiquitous material in semiconductor electronics due to variety
of its applications, its high quality and low prise. Recent successful advancement of nanotechnology
offers routes to develop new concepts in order to create artificial materials possessing novel electronic
and optical properties. Quantum dots (QDs) are promising objects for this goal because of their special
electronic properties as a result of quantum effects due to their small size. Moreover, if QDs are
closely aligned, it results in electron coupling between neighbouring QDs due to overlapping electron
wave functions. Thus, the energy structure of small-period QDs arrays is predicted to be significantly
modified by the artificial periodicity.
Due to 4.2% lattice constant mismatch between Si and Ge, high quality self-assembled Ge QDs can be
grown via the Stranski-Krastanow growth mode by solid source molecular-beam epitaxy (MBE). The
main drawbacks of self-assembled Ge QDs are arbitrary positions where the QDs nucleate as well as
broad size dispersion. In order to circumvent this problem, prepatterned Si substrates are used to
define the position and the size of the QDs.
Fabrication of prepatterned small-period substrates of densely set holes is quite challenging because of
proximity effects and closeness of hole size to resolution limits of available methods. Check-patterns
with different periods down to 35 nm and depth of 5-20 nm are realized by extreme ultraviolet
interference lithography (EUV-IL) and independently by electron beam lithography followed by
subsequent reactive ion etching (RIE) on Si substrates. In this work, influence of both methods on
MBE growth is investigated from the point of view of final quality of the holes and simplicity of
access, usage and precision of positioning on the substrate.
(a)
(b)
Fig 1: AFM images of Ge QDs qrown on (a) flat, (b) prepatterned Si substrate at the same growth parameters.
The prepatterned Si substrates are overgrown first by Si buffer layer, then by few single layers of Ge.
The influence of substrate temperature and buffer layer thickness on homogeneity of Ge dots in order
to optimize growth procedure is investigated by means of atomic force microscopy (AFM). As the
growth takes place at high substrate temperatures, Si tends to reduce strain in Ge via interdiffusion.
This results in formation of effectively SiGe QDs instead of pure Ge. Si distribution is
unhomogeneous in QDs [2] and crucial for artificial energy structure calculation [1]. Nucleation of
QDs in case of small-period arrays is not well studied until now. In-situ UHV scanning tunnelling
microscopy (STM) allows to investigate hole’s shape after buffer growth, morphology of Ge wetting
layer and QDs’ nucleation on the atomic level and to compare them to the unpatterned substrates.
[1] D.Grützmacher, Th. Frommherz, et al., Nanoletters, 7, 3150-3156 (2007).
[2] C. Mocuta, J. Stangl, et al., Phys. Rev. B 77, 245425 (2008).
__________________________
* Contact: s.borisova@fz-juelich.de

We1.5
In situ STM and RHEED study of tensile strained Si grown on
Ge (001) substrates
B. Sanduijav 1* , D. Matei 2 , and G. Springholz 1
1
Institut für Halbleiter- und Festkörperphysik, Johannes Kepler University, A-4040 Linz, Austria
2 Physics of supramolecular systems,University Bielefeld, Universitaetsstr. 2,33615 Bielefeld, Deutschland
Strain engineering in Si/Ge heterostructures is of great importance for tuning the band alignments
in order to obtain effective quantum confinement of holes as well as of electrons in SiGe
heterostructures 1 . Most previous experimental studies have focused on compressively strained SiGe
layers grown on Si substrates. In this case, above a certain critical thickness, strain relaxation proceeds
either by misfit dislocations of by coherent Stranski-Krastanow (SK) islanding, depending on the Ge
content and growth conditions. For confining electrons in SiGe heterostructures, however, tensile
strained layers are required in order to obtain a sufficiently large conduction band offset. Recently,
Pachinger et al. 2 has reported tensile Si growth on Ge substrates, but the resulting Si islands were
found to be dislocated, i.e., plastically relaxed. In our work, we present a detailed in situ scanning
tunneling microscopy (STM) study on the Si island nucleation process as well as their overgrowth
with Ge, combined with real time reflection high energy diffraction (RHEED) studies.
The investigations were carried out in a multi-chamber MBE/STM system equipped with RHEED and
a variable temperature STM. Ge (001) substrates were chemically cleaned similar to Si(001) substrates
using organic solvents and oxygen plasma and a final surface oxide was produced using hydrogen
peroxide for 30s. After introduction in the UHV system, the substrates were outgassed at 500°C
followed by short annealing at 750°C. This results in a flat surface with a typical atomic defects like
missing dimers or adatom clusters. Applying a two-step Ge buffer growth approach produces a large
scale flat surface with a typical RMS of 0.24 nm routinely. Tensile strained Si was then grown on the
Ge buffer at a substrate temperature of 525°C. During the first growth stages, the mixed (1×2) and
c(2×4) typical for Ge remains even up to ~4 ML Si coverage (Fig. 1(a-b)), which transforms to (2×n)-
like reconstruction at higher Si coverages as shown in Fig. 1(b-c). In contrast to (2×n) reconstruction
of missing Ge dimers on Si(001) surface here at this (2×n)-like reconstruction of Si on Ge only every
second Si dimers seem to miss. Moreover, at used growth temperature the Si- wetting layer (WL)
thickness amounts to ~11ML from the STM, where pre-pyramid islands with a density of 0.8µm -2
starts to form as illustrated in Fig. 1(d). This is in good agreement with Pachinger et. al. 2 . Increasing
the coverage to 13.6ML in Fig. 1(e) leads pre-pyramid and pyramid island population with a sidewall
faceting of {113} and a density of 11.1µm -2 . Further increase transforms the Si islands to pre-domes
with additional sidewall faceting of {15 3 23} and a roof faceting of {105} as marked on SOM of
corresponding STM image (Fig. 1(f)) and 3D image of the islands, depicted as insert. In general, all Si
islands have an asymmetric shape, which indicates its relaxation via forming of dislocations as proved
by TEM previously[ 2 ]. To be able to observe dislocation induced surface defects, the Si islands were
capped with Ge at an optimal condition of 350°C growth temperature and a thickness of 50nm. In Fig.
2(a-c) presented STM images are recorded after the Ge cap growth, but RHEED images corresponds
to the end of Si film growth as monitored the growth in real time. After 11ML Si film growth in Fig.
2(a) no dislocation induced defects were formed but only domain boundaries are visible on zoomed-in
picture right to the RHEED pattern. In contrast after 15ML Si growth well developed islands are
formed as recognizable on the RHEED picture and its cap surface contains screw dislocations as
demonstrated in Fig.2(b) and its zoom-in.
This leads to an assumption that the Si WL is not dislocated, but the dislocations form with the island
nucleation. Furthermore only the screw components of the dislocations are clearly remained after
50nm cap layer. As consequence smaller Si coverage of 13.6ML thinner cap layer of 1nm was chosen
that other dislocation components can be visualized. The RHEED pattern corresponding to 13.6ML in
Fig. 2(c) contains lesser intense 3D island spots to compare to 15ML deposition in Fig. 2(b). The
surface topography after capping presents asymmetrically shaped dissolving islands and in addition
line defects, which corresponds to step dislocations as visualized in zoomed in picture.
1 Review in C. Teichert, Phys. Rep. 365, 335 (2002), and references therein.
2 D. Pachinger, H. Groiss, H. Lichtenberger, J. Stangl, G. Hesser, and F. Schäffler, Appl. Phys. Lett. 91, 3606 (2007).

We1.5
11.0ML
3.3ML Si 5.6ML Si 11.0ML
Si
11.0ML Si
13.6ML Si
20.0ML
Si
Pre-pyramid Pyramid Pre-dome
10.7nm
1.3nm
14.8nm
ρ=11.1µm -2 ρ=15.1µm -2
2.7nm
8.6nm
{113} {113}
ρ=0.8µm -2
100nm
21.7nm
21.1nm
100nm
43.6nm
47.1nm
100nm
{15323}
Fig. 1: Ge(001) surface with increasing Si coverage of 3.3-20ML at 525°C. In (a)-(c) a change of surface reconstruction is
visualized at 3.3ML to 11ML, where from 5.6ML a patch (2×n) reconstruction is present. In (d)-(f) Si island evolution is
demonstrated, where pre-pyramids at 11ML and pre-domes at 20ML start to form at this growth temperature. Corresponding
island density is given on STM images and 3D presentation of islands and surface orientation maps (SOM) are depicted as
inserts. Island sizes are noted on the 3D inserts.
11ML Si capped with 50nm Ge
15ML Si capped with 50nm Ge
13.6ML Si capped with 1nm Ge
RHEED 11ML Si Zoom in RHEED 15ML Si
Zoom in RHEED 13.6ML Si Zoom in
Fig. 2: A 50nm Ge cap was grown at an optimized cap growth temperature of 350°C on Si films with thicknesses of (a)
11ML, (b) 15ML. The 11ML film shows on the cap surface no dislocation induced defects in contrast to screw dislocations
visible on 15ML film cap. Capping of 13.6ML Si with only 1nm Ge cap in (c) illustrates a dislocation induced lines. From
real time growth control with RHEED the island formation state prior to the Ge cap growth are demonstrated via
corresponding RHEED pattern, depicted below each STM image. In addition zoomed-in images of surface defects after cap
growth is presented. Large scale image size is 500×500nm 2 .

Wednesday Session We2
Devices

We2.1
MBE growth of IV-VI quantum dots for MIR devices
M. Eibelhuber 1∗ , A. Hochreiner 1 , T. Schwarzl 1 , H. Groiss 1 , W. Heiss 1 , G. Springholz 1
V. Kolkovsky 2 , G. Karczewski 2 , and T. Wojtowicz 2 ,
1 Institute of Semiconductor Physics, University of Linz, Austria
2 Polish Academy of Sciences, Warszawa, Poland
Lead salts have long been used for mid-infrared (MIR) light sources. In particular, the
wide wavelength tunability of IV-VI lasers make them well suited for spectroscopy applications.
The symmetric band structure and small non-radiative Auger recombination
rates provide excellent perspectives for realization of long wavelength lasers with high
operation temperatures. For optically pumped microdisk lasers with strong vertical and
lateral optical confinement, we have recently demonstrated continuous-wave laser
emission at 4.3µm wavelength up to temperatures as high as 0°C [1]. This represents the
highest cw operation temperature for interband lasers in this wavelength region.
For further improvements, quantum dots (QD) as active regions are highly desirable.
However, conventional IV-VI Stranski-Krastanow QDs only exhibit weak luminescence
due to strain-induced type-II band alignment [2]. As an alternative, we have developed a
novel fabrication method for epitaxial PbTe QDs embedded in CdTe matrices [3]. CdTe
and PbTe are practically lattice-matched but exhibit a different lattice structure.
Therefore, these quantum dots are produced by phase separation and interface minimization
rather than by lattice-mismatch strain. As shown by Fig. 1, the resulting QDs
exhibit almost spherical shapes with abrupt interfaces as well as intense
continuous-wave photo-luminescence (PL) at room temperature [3]. The emission
wavelength of the dots can be tuned by adjusting the dot size, which can be achieved by
either changing the amount of deposited PbTe [3] or by varying the growth temperature.
In this work, the structural and optical properties of this unique QD system are
described in detail and the first MIR device applications are presented. In particular, we
show cw QD light emitting diodes (LEDs) operating up to 300 K as well as cw optically
pumped QD lasing up to 200 K. For the latter, PbTe/CdTe microdisks were fabricated
by photolithography and wet chemical etching. A SEM image of a single microdisk is
shown in Fig. 2(d). The QD microdisks were optically pumped at 1030 nm below the
CdTe band gap, resulting in laser emission at around 3.7 µm wavelength up to
temperatures as high as 200 K (Fig. 2(a)). The laser intensity at 200 K as a function of
the pump power is depicted in Fig. 2(b) and exhibits a clear laser threshold. This not
only represents the first quantum dot laser emitting at wavelengths longer than 2 µm but
also further improvements in device structure and processing seem to make room
temperature operation feasible.
[1] M.Eibelhuber et al. Appl. Phys. Lett. 97, 061103 (2010)
[2] M. Simma, et al., Appl. Phys. Lett. 88, 201105 (2006)
[3] W. Heiss, et al., Appl. Phys. Lett. 88, 192109 (2006); H. Groiss, et al., APL 91, 222106 (2007)
∗ Corresponding author: email: martin.eibelhuber@jku.at

We2.1
Fig. 1: (a) Cross-sectional TEM images illustrating the formation process of PbTe quantum dots
in CdTe by interface minimization. In our growth scheme, a low-temperature grown 2D PbTe
layer in embedded in CdTe is subjected to a post growth annealing, which transforms the layer
into highly symmetric isolated PbTe quantum dots without a connecting wetting layer. Thus, the
formation of QDs is not driven by strain but by the immiscibility of layer materials. (b) The
quantum dots show the shape of small rhombo-cubo-octahedrons with atomically sharp
interfaces as proven by the high resolution TEM image on the lower right hand side.
(d)
Fig. 2: (a) Emission spectrum of a PbTe QD microdisk laser and (b) laser intensity versus
effective pump power at 200 K. (c) Cross-sectional TEM image of a PbTe/CdTe reference
sample as in the active region. (d) SEM image of a CdTe microdisk structure with 40 µm
diameter and an active PbTe dot layer in the center of the CdTe waveguide. The undercut in the
GaAs substrate is achieved by selective isotropic wet chemical etching and provides an effective
mode confinement.

We2.2
The role of doping scheme in ultra-low disorder MBE grown
mesoscopic FQHE devices
V. Umansky* , M. Heiblum, M. Dolev and Y. Gross
Braun Center for Submicron Research, Weizmann Institute of Science, Rehovot 76100, Israel
Nowadays an ongoing research in the field of Fractional Quantum Hall Effect (FQHE) focuses on
exotic fractional states, like the 5/2 state with predicted non-Abelian statistics, as well as on recently
experimentally discovered new types of quasiparticles that do not carry any charge, but only energy
along the sample edge - so called neutral edge modes.
A vast majority of these studies are carried out using 2DEG in AlGaAs/GaAs. Despite of a common
approach that only very pure 2DEG exhibits clear FQHE states, especially the fragile 5/2 state, it has
been recently shown [1] that low background impurity density in GaAs and, hence, very high mobility,
by itself, is not sufficient to produce samples with high quality FQHE states. Potential fluctuations
introduced by remote ionized impurities were found to play a crucial role in defining the energy gap of
the quasiparticles. Moreover, comparison of recently measured energy gap of the 5/2 state with
theoretical values suggests that the qusiparticles energy gap is very sensitive to small topological
details of the disorder landscape, thus the FQHE quality strongly depends on the spatial distribution of
the scattering centers (and not only on their density) [2]. Since in MBE growth the doping atoms are
introduced in a completely random manner, the only way to facilitate “disorder engineering” is to
create Coulomb correlation among carriers and ionized impurities in the doping layers.
The standard doping schemes utilizing Silicon as a donor impurity in the AlGaAs with Aluminum
mole fraction 0.3÷0.4 are not efficient in creating correlated system of carriers in the doping layer due
to formation of deep DX center level(s) at relatively high freeze-out temperature (100÷140K). One of
the methods that potentially allow high degree of correlations
is doping inside Short Period Supper-lattice (SPSL)
suggested in [3] and realized for ultra-high mobility
structures in [1]. Indeed very high quality FQHE was
achieved (Fig.1), however, the fabricated quantum devices
(for instance Quantum Point Contacts) exhibited insufficient
temporal stability due to poor localization of electrons in the
doping layer even at 10 mK temperature and, hence, high
sensitivity to any fluctuation in the external electric field.
Recently, a pronounced hysteresis upon gating has been
described in similar structures [4].
Finding a proper compromise between highly spatial
correlated electron system combined with strong localization
of the carriers in the doping layer, needed for low
temperatures mesoscopic device stability, creates an
unprecedented challenge both for heterostructure design and for the MBE growth process
implementation. We describe the results on application of various doping schemes in ultra-low
disorder 2DEG heterostructures. Promising results were obtained using precisely controlled delta
doping in relatively low Aluminum mole fraction layers containing DX centers. We discuss the
possible physical mechanisms leading to substantially lower disorder and better stability in these
structures.
R
xy
(kΩ)
R
xx
(kΩ)
12
8
4
0
0.6
0.4
0.2
0.0
T=10 mK
0 1 2 3 4 5 6 7
B (T)
n s
= 3.1·10 11 cm -2
µ = 30·10 6 cm 2 /V·s
Fig 1: Hall and longitudinal resistance
in SPSL doped sample
3
5
2
2
[1] V. Umansky, M. Heiblum, Y. Levinson, J. Smet, J. Nübler , M. Dolev, J. Crystal Growth, 311 ,1658,(2009)
[2] J. Nuebler, V. Umansky, R. Morf, M. Heiblum, K. von Klitzing, J. Smet, Phys. Rev. B, 81, 035316 (2010)
[3] K.J. Friedland, R. Hey, H. Kostial, R. Klann, K. Ploog, Phys. Rev. Lett. 77, 4616 (1996).
[4] C. Roessler, T. Feil, P. Mensch, T. Ihn, K. Ensslin, D. Schuh, W. Wegscheider, New Journal of Physics, 12, 043007
(2010)
__________________________
* Contact: Vladimir.Umansky@weizmann.ac.il

We2.3
Investigations of Si-dopant layers on ultrahigh-mobility 2DEGs
in GaAs/AlGaAs-structures
C. Reichl 1* , E. de Wiljes 1 , C. Rössler 1 and W. Wegscheider 1
1 ETH Zürich, Laboratorium für Festkörperphysik, 8093 Zürich, Switzerland
Two-dimensional electron gases (2DEGs) in Al x Ga x-1 As-heterostructures with ultra-high mobilities
exceeding 10 7 cm 2 V -1 s -1 are of great interest for different research fields in solid states physics. The
prominent fractional quantized Hall state at ν = 5/2, microwave-induced zero-resistance states or
interactions between composite fermions can be observed only in high-quality samples displaying
such mobilities [1]. Essential for further studies on these phenomena is the optimization of the 2DEG’s
electron density.
Here we report on the effects of δ-Si doping layers in AlAs-GaAs-AlAs quantum wells, providing
ultra-high mobilities as well as 2DEG densities insensitive to illumination. Carefully designed
thicknesses of these AlAs- and GaAs-layers provide X-states in AlAs slightly lower in energy than the
Γ-states in the GaAs-layer. This provides an effective screening of the 2DEG from the disordered
remote ionized (RI) impurity potential [2, 3].
In addition it eliminates DX-centers for roughly doubling electron densities with no higher RIscattering,
thus making illumination of the sample unnecessary. Samples were grown with mobilities
>16·10 6 cm 2 V -1 s -1 with a density of ~3·10 11 cm -2 at 1,5 Kelvin without illumination.
Furthermore, the screening effects of the AlAs barrier allow for a wide range of doping density at no
cost for mobility. Doping was varied from 1,1·10 16 cm -2 to 4·10 16 cm -2 with no effect on 2DEG-density
(3,0·10 11 cm -2 ) and mobility (16,0·10 6 cm 2 V -1 s -1 to 16,2·10 6 cm 2 V -1 s -1 ).
This wide range of doping is also an important prerequisite for further work on 2DEG-based Qbits,
where electron density is varied by Schottky-gates. It was shown that top-gated structures can be used
to tune the density, however, pronounced gate voltage hysteresis accompanied by a degradation of the
mobility occurs [4]. Influencing both doping layers independent from each other would be desirable.
Hence, high-mobility samples with a burried backgate, insulated by a layer of low-temperature GaAs,
were grown. In such structures, the electron density can be tuned via an applied backgate voltage by
around 20%, as it was done with top-gate voltage.
Fig 1: Mobility (squares) and density (dots) plotted against
thickness of the GaAs quantum well (AlAs-layers
fixed at 2 nm).
Fig 2: Electron density (upper) and mobility (lower graph)
versus time. Backgate voltage was +1V for 300s, then
switched to -1V for 180s and finally to -2V.
[1] Pfeiffer L and West K W, 2003, Phys. E 20 57
[2] Friedland K J et al., 1996, Phys. Rev. Let. 77, 224616
[3] Umansky V et al., 2009, J. of Crystal Growth 311, 1658
[4] Rößler C et al., 2010, New J. Phys 043007
__________________________
* Contact: creichl@phys.ethz.ch

We2.4
Short wavelength high power Quantum Cascade Lasers
X. Marcadet 1 , M. Carras 1 , B. Simozrag 1 , M. Garcia 1 , G. M. De Naurois 1 ,
G. Maisons 1 , O. Parillaud 1 , O. Patard 1 , F. Pommereau 1 , O. Drisse 1 , F. Alexandre 1 ,
J. Massies 2
1
Alcatel Thales III-V Lab, 1 avenue Augustin Fresnel, 91767 Palaiseau cedex, France
2 CRHEA-CNRS, 06560 Valbonne, France
The need of high power sources in the mid IR for defense application has raised much effort
recently to bring Quantum Cascade Lasers (QCLs) to a new level of performances. Sources above
Watt level have been demonstrated thanks to the mastering of thermal management [1-2]. Our work
was to provide a reliable, easy to integrate high power QCL source. The QCL active region are grown
by molecular beam epitaxy on 2 inches substrates using a Riber 49 multi-wafer production machine.
To obtain QCL emission below 5 μm at room temperature (RT) with the GaInAs/AlInAs material
system epitaxially grown on InP substrates, it is necessary to maximise the conduction band offset of
the heterostructure (ΔEc) to improve the device thermal behavior. This implies in turn that highly
strained barrier layers have to be used. The strain balance should be carefully controlled, a point which
remains challenging. High-resolution X-ray diffraction of double Ga x In 1-x As/Al y In 1-y As superlattices is
used to calibrate the composition and growth rates of the strain-balanced layers. This approach is
shown to be an effective way to extract from a single HRXRD experiment the complete set of the
structural parameters defining the SLs.
In a first time we have developed sources around 1 watt at room temperature in continuous
wave mode with a standard technology using a dielectric coating of the laser ridge. We have obtained
TM00 sources above 500mW at 4.6µm (figure 1) with a reproducibility of less than 1 percent from
laser to laser and decent wall plug efficiency of 6.5%. TM00 sources above 200 mW have been
obtained at 3.9 µm. We will discuss on the effect of the increase of the barrier height while keeping the
same emission wavelength on the performances of the lasers. With this technology and at these
wavelengths, beam quality degrades when power further increases.
In a second part we present the development of semi-insulating Fe-doped buried
heterostructures. The active regions being the same as for standard technologies, we can compare
advantages and drawbacks of buried lasers. Characterization of buried laser show that thermal
resistance can be easily improved by at least a factor 2. Current leakage at the mesa edges remains a
major issue to be solved to insure reliable and reproducible results using buried heterostructures.
Results at 4.6 µm and 3.9 µm are shown for different strategies in the preparation of the edged mesa
before InP:Fe regrowth. Figure 2.a shows a typical buried structure profile. Compared performances
between standard and buried technologies are shown in figure 2.b.
In a third part we will discuss on the ways to further improve the power and wall plug efficiency
without degrading important values like the beam quality. An option is beam combining whereby
several QCLs are phase-locked to create an extraordinarily high power beam with a single transverse
mode.
[1] Y. Bai, S. Slivken, S.R. Darvish, and M. Razeghi, “Very high wall plug efficiency of quantum cascade lasers” SPIE
Proceedings, San Francisco, CA (January 22-28, 2010), Vol. 7608, p. 76080F-1-- January 22, 2010
[2] A. Lyakh, R. Maulini, A. Tsekoun,1 R. Go, C. Pflügl, L. Diehl, Q. J. Wang, Federico Capasso, and C. Kumar N. Patel, “3
W continuous-wave room temperature single-facet emission from quantum cascade lasers based on nonresonant extraction
design approach”, Appl. Phys. Lett. 95, 141113 (2009)
_________________________
* Contact: b.author@institution1.edu

Voltage (V)
Output power (mW)
Detected power (a.u)
Voltage (V)
Optical power (mW)
We2.4
12
10
8
6
4
HR coating
Epi down
400
L=4 mm
2
CW
200
RT on TEC
0
0
0.0 0.5 1.0 1.5 2.0
Current (A)
(a)
14 µm
12 µm
10 µm
8 µm
6 µm
1000
800
600
0.05
0.04
0.03
0.02
0.01
0.00
-40 -20 0 20 40 60
Angle (°)
Fig 1: (a) I(V) and P(I) characteristics of a 4.6 µm QCL with a standard technology process, (b) corresponding
Farfield for a 6µm ridge width.
(b)
14 Buried
12
Standard
300
10
200
(a)
8
6
4
2
0
HR coating
Epi down
L=3 mm
l= 6 µm
CW
RT on TEC
0
0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7
Current (A)
Fig 2: (a) SEM picture of a detail of the regrown InP:Fe along the GaInAs/AlInAs QCL active region. (b)
Comparison of I(V) and P(I) curves between standard and buried 4.6 µm QCLs with the same active region.
(b)
100

We2.5
MBE growth of InGaAs/GaAsSb based
mid-infrared and THz quantum cascade lasers
H. Detz 1,* , A.M. Andrews 1 , P. Klang 1 , C. Deutsch 2 , M. Nobile 1 ,
W. Schrenk 1 , K. Unterrainer 2 and G. Strasser 1
1
Center for Micro- and Nanostructures and Institute for Solid-State Electronics,
Vienna University of Technology, Floragasse 7, 1040 Wien,Austria
2
Photonics Institute, Vienna University of Technology
Quantum cascade lasers are well established coherent light sources for the mid-infrared (MIR) [1]
and THz [2] spectral regions. Each cascade is a sequence of quantum wells and barriers, where
electrons can relax in a radiative transition, which can be repeated for higher efficiency. One approach
to improve their performance and efficiency is the use of novel material combinations. Recently QCLs
based on InGaAs/GaAsSb heterostructures were demonstrated as a suitable alternative for
GaAs/AlGaAs with a comparable conduction band offset of 0.36 eV [3]. An important advantage,
compared to the GaAs/AlGaAs or the InGaAs/InAlAs material system, is the low effective mass both
in the InGaAs wells and GaAsSb barrier layers, allowing higher optical matrix elements for increased
gain and thicker barriers, which are less sensitive to interface roughness or deviations in the growth
rate [4].
The In 0.53 Ga 0.47 As/GaAs 0.51 Sb 0.49 heterostructures were grown lattice-matched to InP substrates in a
solid-source MBE system with valved crackers for both group V species. After desorbing the native
oxide under an As 4 flux, the wafers were brought to a growth temperature of 480°C, which was
maintained throughout the whole growth process. Typical InGaAs and GaAsSb growth rates were 1
and 0.5 µm/h respectively. Correct growth rates and compositions as well as the overall quality were
ensured by X-ray diffraction measurements around the 004, 002 and 224 diffraction peaks [5]. Based
on the first InGaAs/GaAsSb MIR QCL [3], an improved 4-well active region with double phonon
depletion for emission wavelengths around 11 µm was designed. Sixty active periods, sandwiched
between low loss plasmon-enhanced InGaAs waveguide layers were grown on higly doped InP
substrates. Fabry-Pérot resonator QCLs processed as 50 µm ridge waveguides reached a maximum
optical output power of 1.2 W at 78 K with a low threshold current density of 0.6 kA/cm².
Furthermore, the first InGaAs/GaAsSb based THz QCLs were realized [6]. The structure consisted of
170 3-well active periods for a total thickness of 10.2 µm. The minimum barrier thickness of 1 nm,
which is three times larger than for InGaAs/InAlAs based THz QCLs [7], allows a reproducible layer
thickness due to reduced effects from substrate rotation and shutter transients for this long growth
time. These devices showed emission around 80 µm, which corresponds to 3.6 THz with a maximum
operation temperature of 135 K, already surpassing the InGaAs/InAlAs material system.
In summary, we demonstrated InGaAs/GaAsSb based QCLs as an approach for more efficient
intersubband device designs, covering the wavelength regions around 11 and 80 µm. The low effective
mass in the barriers should allow improved gain and therefore higher operating temperatures of THz
QCLs.
[1] J. Faist, F. Capasso, D.L. Sivco, C. Sirtori, A.L. Hutchinson and A.Y. Cho, Science 264, 553 (1994)
[2] R. Köhler, A. Tredicucci, F. Beltram, H.E. Beere, E.H. Linfield, A.G. Davies, D.A. Ritchie, R.C. Iotti and F. Rossi,
Nature 417, 156 (2002)
[3] M. Nobile, P. Klang, E. Mujagić, H. Detz, A.M. Andrews, W. Schrenk and G. Strasser, Electron. Lett. 45, 1031
(2009)
[4] E. Benveniste, A. Vasanelli, A. Delteil, J. Devenson, R. Teissier, A. Baranov, A.M. Andrews, G. Strasser, I. Sagnes
and C. Sirtori, Appl. Phys. Lett. 93, 131108 (2008)
[5] H. Detz, A.M. Andrews, M. Nobile, P. Klang, E. Mujagić, G. Hesser, W. Schrenk, F. Schäffler and G. Strasser, J.
Vac. Sci. Technol. B 28, C3G19 (2010)
[6] C. Deutsch, A. Benz, H. Detz, P. Klang, M. Nobile, A.M. Andrews, W. Schrenk, T. Kubis, P. Vogl and G. Strasser,
to be published in Appl. Phys. Lett.
[7] M. Fischer, G. Scalari, C. Walther and J. Faist, J. Cryst. Growth 311, 1939 (2009)
__________________________
*
Contact: hermann.detz@tuwien.ac.at

We2.5
Fig 1: Light output vs. current density of a 50 µm wide InGaAs/GaAsSb MIR QCL. The inset shows
a typical Fabry-Pérot emission spectrum of the same device. The measurements were done at 78 K.
Fig 2: Light output vs. current density of a 40 µm wide double metal waveguide InGaAs/GaAsSb
THz QCL. Emission spectra for two current levels are plotted in the inset. All data were recorded at
7 K.

We2.6
Room temperature operation of a GaInAsSb/AlGaInAsSb digital
alloy laser diode at 3.3 µm
S. Belahsene 1* , K. S.Gadedjisso 1 , G. Boissier 1 , P. Grech 1 , G. Narcy 1 and Y.
Rouillard 1
1Institut d'Electronique du Sud, UMR 5214 CNRS, Université Montpellier 2, Montpellier34095, France 1
During the past decade, remarkable progress has been made in the development of room-temperature
operating GaInAsSb/AlGaInAsSb lasers emitting above 3µm [1-4]. The introduction of quinary
AlGaInAsSb barriers is an alternative way to improve hole confinement in the GaInAsSb quantum
wells (QWs). The growth of this material is delicate due to the high content of arsenic in it (y As =
0.24). A relative uncertainty of 3 % on this content is necessary to keep sufficient lattice matching.
This uncertainty actually corresponds to a limit of what can be done by molecular epitaxy. To
overcome this problem, we have used an alternative material, the so called “digital alloy” [5]. The key
point here is to separate the arsenic and antimony containing materials in order to make the
incorporation of arsenic independent of the arsenic and antimony fluxes and of the growth
temperature.
The laser structure was grown by molecular beam epitaxy (MBE) in a Varian Gen II system equipped
with two valved As and Sb cracker cells. The growth was carried out on an n-type (100) GaSb
substrate. It is constituted of an active region containing four 11 nm-thick compressively strained QWs
with an average composition of Ga 0.45 In 0.55 As 0.20 Sb 0.80 . The quantum wells were constituted of the
following digital alloy: (InAs) 1.2 ML (Ga 0.57 In 0.43 Sb) 4.6 ML (the period of the pattern is below ~8 MLs [5])
embedded between Al 0,25 Ga 0,51 In 0,24 As 0.24 Sb 0.76 barriers constituted of the digital alloy: (InAs) 2
ML (Ga 0.67 Al 0.33 Sb) 6 ML . The growth sequence is shown in Figure 1. Growth interruptions under arsenic
or antimony flux were used to adjust the lattice mismatch of these materials.
High Resolution X-Ray diffraction around the (0 0 4) substrate reflection was used to determine the
strain of the QWs and the lattice mismatch of the barrier which are respectively of 1.4% and 0.56 10 -3
(Figure 2). Figure 3 shows the band structure of the active zone. The calculated value of the offset in
the valence band equals 180 meV providing a good confinement for holes.
Broad area laser diodes were made from the grown wafer. Figure 4 shows the power-current
characteristic for an uncoated 100 x 2300 µm 2 device obtained at RT in the pulsed regime (1%, 1µs),
the threshold current density is 1094 A/cm 2 and the differential efficiency η d is 14 %. The spectrum
measured at room temperature is shown in Figure 5. The emission wavelength is 3.3 µm
corresponding to a strong absorption line of CH 4 .
[1] M. Grau, C. Lin, O. Dier, C. Lauer, and M.-C. Amann, Appl. Phys. Lett. 87, 241104 (2005).
[2] T. Hosoda, G. Belenky, L. Shterengas, G. Kipshidze, and M. V. Kisin, Appl. Phys. Lett.92, 091106 (2008).
[3] L. Shterengas, G. Belenky, T. Hosoda, G. Kipshidze, and S. Suchalkin, Appl. Phys. Lett.93, 011103 (2008).
[4] T. Hosoda, G. Kipshidze, L. Shterengas, S. Suchalkin, and G. Belenky, Appl. Phys. Lett.94, 261104 (2009).
[5] Ron Kaspi and Giovanni P. Donati, J. Crystal Growth 251 (2003) 515-520.
__________________________
* Contact: belahsene@ies.univ-montp2.

Barrier
Al 0,25Ga 0,51In 0,24As 0.24Sb 0.76
Spacer
Al 0,25Ga 0,51In 0,24As 0.24Sb 0.76
QWs
GaSbBe
Al0,10Ga0,90As0.07Sb0.93 ( 1.5 µm)
GaAlSb(6ML)
InAs(2ML)
GaAlSb(6 ML)
InAs(2ML)
GaInSb(4.6ML)
InAs(1.2ML)
Ga 0.45In 0.55As 0.20Sb 0.80 -500
GaAlSb(6ML)
InAs (2ML)
Al0,10Ga0,90As0.07Sb0.93 ( 1.5 µm)
Substrat: GaSbTe
x40
x12
x6
x40
x4
Energie (meV)
1500
1000
500
0
GaSb
InAsSb
Epaisseur (Α 0 )
We2.6
E g
= 370 meV
Fig 2: Band structure of the active zone with an
InAs/GaInSb well and InAs/GaInSb barriers
E g
180 meV
Fig 1: Schematic diagram of the structure
1.E+06
1.E+05
A 3 0 3
Al0.25Ga0.51In0.24As0.22Sb0.78
ℵa/a = +0.56 x 10-3
GaSb
Intensity (cts.s-1)
1.E+04
1.E+03
1.E+02
Ga0.45In0.55As0.29Sb0.71
ℵa/a = +1.40 x 10-2
1.E+01
1.E+00
28.80 29.30 29.80 30.30 30.80
Omeg a ( °)
Fig 3: High resolution X-Ray scan around the (004) reflection
A 3 0 3 - 2 3 0 0 - 0 1
P (mW)
180
160
140
120
100
80
60
294 K, 1 %
ith = 2516 mA
Jth = 1094 A/cm²
nd = 14 %
Pmax = 153 mW
PI-V1297-2300-03
Intensity (a.u.)
1.4
1.2
1
0.8
0.6
0.4
21°C, D.C. = 1 % , P.W. = 0.2 µs
Lambda Peak = 3.32 µm
40
0.2
20
0
0 2000 4000 6000 8000 10000
i (mA)
0
2.9 3 3.1 3.2 3.3 3.4 3.5 3.6 3.7
Wavelength (µm)
Fig 4: P (I) characteristic of a 100 µm-wide
2300-µm long broad area laser diode
Fig 5: Spectrum of a 100 µm-wide 2300-
µm long broad area laser diode

Wednesday Session We3
Nanowires

We3.1
Growth kinetics of III-V nanowires
F. Glas*, J. C. Harmand and G. Patriarche
CNRS - Laboratoire de Photonique et de Nanostructures, Route de Nozay, 91460 Marcoussis, France
The growth of nanowires in the vapor-liquid-solid (VLS) mode, whether it be catalyzed by foreign
particles or self-catalyzed, involves the nucleation of new solid monolayers from a supersaturated
liquid nanodrop. In a stationary or quasi-stationary regime, the material consumed by forming the
nanowire is balanced by that fed to the drop from the vapor. The latter has three main components,
resulting respectively from direct impingement on the nanodrop, collection from the substrate and
collection by the sidewalls, the latter two followed by surface diffusion to the drop.
After briefly introducing this general framework, we will show how the combination of growth
experiments, transmission electron microscopy (TEM) and growth modeling permits to analyze
quantitatively some of the relevant nucleation and growth processes. While emphasizing the general
characters of VLS growth, we shall also point out some specificities of MBE and of nanowires of
compound semiconductors.
We developed a method for measuring the growth chronology of individual nanowires that consists in
establishing a weak time-periodic composition modulation along the nanowire axis and measuring by
ex situ TEM the lengths grown during each period [1]. In this way, the instantaneous growth rate of the
nanowire becomes accessible. By fitting the experimental data with models, we derive values of the
adatom diffusion lengths and of the chemical potentials in the adsorbed and liquid phases. We shall
discuss the relative importance of the three abovementioned material contributions to nanowire
growth.
The VLS growth of nanowires offers a privileged example of nucleation in an open system of
nanometric dimensions. The addition of each solid monolayer to the nanowire indeed requires the
formation of a new 2D nucleus at the interface between the nanowire and the liquid nanodrop, most
often at the triple VLS line [2]. Classical nucleation theory assumes that the consecutive nucleation
events are temporally independent. We show that this postulate, although reasonable for a macrophase,
must be abandoned for a nanophase. By using the abovementioned method, we count the numbers of
nucleations occurring during consecutive and equal time intervals. In the case of Au-seeded InP x As 1-x
nanowires, the corresponding statistics are markedly sub-Poissonian, which proves that the nucleation
events are temporally anticorrelated. This nucleation antibunching stems from the depletion of the
nanodrop and the decrease of chemical potential that follow the formation of each monolayer [3]. The
effect is all the more marked that the drop is small and contains a low concentration of nanowire
atoms. Such is indeed the case for group V species during the growth of III-V nanowires.
We have developed a self-consistent growth model which accounts for the impact of these variations
of concentration and of the associated changes of chemical potential on nucleation. In the simulations,
we use original chemical potential calculations allowing for the fact that the drop is a ternary Au-III-V
liquid [4]. These calculations also open the way to an improvement of the current nanowire growth
models by fully taking into account the two components of compound semiconductors. Moreover, we
previously argued that the wurtzite/sphalerite polytypism observed in nanowires of cubic III-V
compounds is partly governed by the value of the chemical potential in the nanodrop [2]. In this
context, we briefly discuss the implications of nucleation antibunching on the occurrence of these two
crystal structures.
[1] J. C. Harmand, F. Glas and G. Patriarche, Phys. Rev. B 81, 235436 (2010).
[2] F. Glas, J. C. Harmand and G. Patriarche, Phys. Rev. Lett. 99, 146101 (2007).
[3] F. Glas, J. C. Harmand and G. Patriarche, Phys. Rev. Lett. 104, 135501 (2010).
[4] F. Glas, J. Appl. Phys. 108, 073506 (2010).
__________________________
* Contact: frank.glas@lpn.cnrs.fr

We3.2
InAs Quantum Dot Arrays Decorating the Facets of GaAs Nanowires
E. Uccelli 1,2,* , J. Arbiol 3 , J.R. Morante 4 and A. Fontcuberta i Morral 1,2
1
Laboratoire des Matériaux Semiconducteurs, Ecole Polytechnique Fédérale de Lausanne, Switzerland
2
Walter Schottky Institut and Physics Department, Technische Universität München, Germany
3
Institució Catalana de Recerca i Estudis Avançats and Institut de Ciència de Materials de Barcelona, Spain
4
Departament d’Electrònica, Universitat de Barcelona and Catalonia Institute for Energy Research, Spain
Semiconductor Nanowires (NWR) are being widely investigated, motivated by a desire to discover
and manipulate physics at nanometer dimensions as well as by being considered promising building
blocks in future applications such as energy harvesting or biosensing. Heterostructures in the axial as
well as radial direction are examples for novel device structures [1,2]. In this work, we present a new
approach to define optically active Stranski-Krastanov InAs Quantum Dot (QD) on the sidefacets of a
GaAs NWR [3].
Catalyst-free GaAs nanowires have been grown via the Ga assisted MBE method [4,5]. The nanowires
form along the [111] direction and present a hexagonal prism morphology, with the facets pertaining to
the {110} family. The deposition of InAs quantum dot chain on (110) surfaces has been a challenging
task, because InAs quantum dots tend to form on (001) surfaces but not on (110) facets. However, we
were able to show that highly ordered InAs QD arrays can be realized on (110) AlAs oriented surface.
Indeed, after the formation of a GaAs nanowire, we have first covered the nanowire facets with a
coaxial thin AlAs shell layer, and then deposited a few monolayers of InAs. We have found that InAs
nucleates forming a QD chain on the side facet of the nanowire.
The morphology and structure of these heterostructures have been investigated by atomic force and
high resolution transmission electron microscopy, together with energy-dispersive X-ray spectroscopy
and electron energy loss spectroscopy. Single and dual quantum dot chain structures have been
observed, with a typical QD height in the range of 10 nm. The optical properties have been
characterized by low temperature micro-photoluminescence, reporting highly intense features between
1.35 eV and 1.42 eV. Sharp peaks have been measured on the single nanowires and correspond to the
radiative recombination of single and double excitons in the quantum dots. The emission energy
depends on the size of the QDs. This new type of 1D-0D combined heterostructures can open a new
avenue to the fabrication of highly efficient single-photon sources, novel quantum optics experiments,
as well as the realization of intermediate-band nanowire solar cells for third-generation photovoltaics.
Fig 1: (a) Schematic drawing of the sample. (b) Power dependent photoluminescence series on InAs QD.
[1] L. Samuelson et al., Physica E 25, 313 (2004)
[2] A. Fontcuberta i Morral et al., Small, 4, 899 (2008)
[3] E. Uccelli et al., ACS Nano 4, 5985 (2010)
[4] C. Colombo et al., Phys. Rev. B 77, 155326 (2008)
[5] A. Fontcuberta i Morral et al., Appl. Phys. Lett. 92, 063112 (2008)
__________________________
* Contact: emanuele.uccelli@epfl.ch

We3.3
AlAs-GaAs core-shell nanowires
grown by chemical beam epitaxy
A.LI 1,* , D. Ercolani 1 , F. Rossi 2 , L. Nasi 2 , G. Salviati 2 , F. Beltram 1 and L. Sorba 1
1
NEST, Istituto Nanoscienze-CNR and Scuola Normale Superiore, I-56127 Pisa, Italy
2 IMEM-CNR, Parco Area delle Scienze 37/A, I-43100, Parma, Italy
Despite the technological importance of AlAs and AlGaAs alloys, so far exploitation of AlAs in
nanowires (NWs) was limited to thin layers in heterostructures with GaAs. To our knowledge, neither
pure AlAs NWs nor AlAs-GaAs core-shell NWs were reported, so that as yet very little is known on
crystal structure or electronic and optical properties of these nanostructures.
AlAs NWs and AlAs-GaAs core-shell NWs were grown on GaAs 111 (B) substrates by Au-assisted
chemical beam epitaxy. For the growth of AlAs NWs, TMAl and TBA were used as precursors, and
TEG and TBA for the GaAs shells. NWs were investigated in situ by reflection high-energy electron
diffraction (RHEED) and ex situ by scanning electron microscopy (SEM), high resolution transmission
electron microscopy (HRTEM) with energy dispersive X-ray spectrometry (EDS) and scanning
transmission electron microscopy (STEM). In situ RHEED indicates that AlAs NWs have wurtzite
crystal structure. Pure AlAs NWs readily oxidize when they are exposed to air, as shown by TEM and
EDS experiments.
We optimized the growth of a thin (~7 nm thick) GaAs shell around the AlAs NWs following AlAs
growth. A two-step temperature procedure was employed to grow the GaAs shell. Not only lateral but
also axial growths of GaAs were observed by ex situ STEM (see Fig. 1). We studied the relationship
between axial growth and growth temperature. We found that axial GaAs growth can be controlled
without substantial change in shell thickness (see Fig. 2). Structure characterization of AlAs-GaAs
core-shell NWs was performed in situ by RHEED and ex situ by SEM, TEM and EDS. Results show that
AlAs-GaAs core-shell NWs have wurtzite crystal structure and that the GaAs shell can successfully
prevent rapid oxidation of the AlAs core.
60
c)
GaAs axial growth rate (nm/min)
50
40
30
20
10
a)
b)
10 nm
0
320 340 360 380 400 420 440 460 480 500
GaAs growth temperature ( O C)
Fig 1: (a) STEM image of AlAs-GaAs NWs,
(b) HRTEM image of AlAs-GaAs NWs,
(c) Fast Fourier Transform analysis.
Fig 2: GaAs axial growth rate vs. GaAs
growth temperature.
__________________________
* Contact: ang.li@sns.it

We3.4
Distinct nucleation and growth modes of self-assisted InAs
nanowires on bare Si(111)
E. Dimakis * , J. Lähnemann, U. Jahn, S. Breuer, M. Hilse, L. Geelhaar, and H.
Riechert
Paul Drude Institute for Solid State Electronics, Hausvogteiplatz 5-7,10117 Berlin, Germany
InAs nanowires (NWs) can be epitaxially grown on Si without using any catalysts. Satisfactory results
have been obtained when the substrate is masked by a Si-oxide layer with small holes that have been
spontaneously or lithographically opened; and in this case NW nucleation occurs inside these holes.
Nevertheless, the growth mechanism is not well understood and contradictory reports are found in the
literature. While for self-assisted GaAs NWs the presence of Ga droplets at the NW tips evidences a
Ga-droplet-mediated vapor-liquid-solid (VLS) growth mode, for InAs it is not clear whether the
formation of NWs results from an In-droplet-mediated VLS or a droplet-free vapor-solid (VS) growth
mode. Furthermore, high densities of stacking faults and/or wurtzite inclusions in predominantly zinc
blende NWs are typically reported. Thus, it is necessary to develop a better understanding of the
growth mechanism in order to improve the growth control and the crystal quality of InAs NWs on Si.
Here, we present InAs growth studies conducted on bare Si(111) substrates by solid-source molecular
beam epitaxy, without using any catalyst-particles, (native) oxide or other coating layers. In this way,
the growth is investigated without the influence of foreign elements or induced morphological
features, and the conditions leading to NW growth are explored. Each of the growth parameters,
namely the growth temperature, the indium atomic flux, the V/III flux ratio, and the growth duration,
is studied independently in a systematic way.
We show that InAs NWs do form and grow aligned parallel to the substrate surface normal. Hence, the
patterning or masking of Si is not a pre-requisite for the formation of NWs. These NWs are not tapered
and have hexagonal cross sections with well-defined facets. The NW tips end with a flat top facet
parallel to substrate surface, and no In droplets exist. Reflection high energy electron diffraction
(RHEED) and electron backscatter diffraction (EBSD) were used for the structural characterization of
the NWs during and after growth, respectively. We find that the NWs have the wurtzite structure with
{112¯ 0} side facets, in contrast to what is typically reported for self-induced NWs on oxidized Si. The
epitaxial relationship with the substrate is: {112¯ 0}InAs║{11¯ 0}Si, {0001}InAs║{111}Si.
We analyze how the NW number density, diameter, and axial growth rate depend on the growth
conditions and duration, and deduce two distinct mechanisms, one for the nucleation and one for the
growth stage. The nucleation of InAs NWs is associated with the formation of In droplets, while the
subsequent growth proceeds in a VS mode; in the latter stage the conditions at the NW tips are Asrich.
This transition from In-rich to As-rich conditions at the NW tips upon NW nucleation is
correlated with the fast increase of the number of nucleated NWs in the beginning of the growth. The
validity of the suggested mechanisms is confirmed by additional experiments employing a two-step
process, where the growth conditions were intentionally changed after the NW nucleation was
completed.
As shown by our experiments, the formation of In droplets (and thus the number density and the size
of the NWs) is limited by the surface diffusivity of In adatoms on Si, which in turn is determined by
the growth conditions used. However, comparative studies on bare and oxidized Si suggest that the
diffusivity of In adatoms on oxidized Si is mainly limited by morphological features of the oxide
surface. In that case, the growth conditions need to be adjusted to avoid nucleation on the oxide layer.
__________________________
* Contact: dimakis@pdi-berlin.de

We3.4
(b)
1µm
(a)
(c)
Fig 1: (a) Typical cross sectional SEM image of InAs NWs grown on bare Si(111). (b) Top-view SEM image showing
that the NWs have hexagonal cross sections with non-tapered well-defined facets. (c) Side-view SEM image showing
the absence of In-droplets from the flat top facets at the NW tips.
{0001}
_
{1120}
(a)
(b)
(c)
Fig 2: (a) EBSD pattern as measured from a NW side facet under an angle of 70 o . The colored lines correspond to the
fitted Kikuchi pattern of a wurtzite InAs crystal. (b) The fitted Kikuchi pattern as corrected for the angle of 70 o . (c)
Schematic representation of the NW facet crystallographic orientation as suggested by the EBSD analysis.
(5)
(1) (2) (3)
(4)
Fig 3: Schematic representation of the suggested model for the nucleation and growth of InAs NWs on Si(111). (1)-
(2) The nucleation of InAs NWs occurs under In-droplets initially formed on Si. (3)–(5) The excess of In is soon
consumed due to the ongoing nucleation of new NWs, and the growth continues under excess of As, in a VS mode.

We3.5
Strain balanced technique for the growth of very high aspect ratio
quantum posts
D. Alonso-Álvarez 1, *, B. Alén 1 , J. M. Ripalda 1 , J. Llorens 1 , A. G. Taboada 1 , Y. González 1 ,
L.González 1 , F.Briones 1 , M. A.Roldán 2 , J.Hernandez-Saz 2 , M.Herrera 2 and S.I.Molina 2
1 IMM-Instituto de Microelectrónica de Madrid (CNM-CSIC), Isaac Newton 8, 28760 Tres Cantos, Spain
2 Departamento de Ciencia de los Materiales e Ing. Metalúrgica y Q. I. Universidad de Cádiz, Campus
Universitario de Puerto Real, 11510 Puerto Real, Cádiz, Spain
The versatility of the molecular beam epitaxy (MBE) has encouraged many groups to explore
new kind of nanostructures beyond quantum wells, wires and dots (QDs) with exciting and useful
properties. That is the case of quantum posts (QPs) which are the limit of columnar QDs when the
spacer between layers is reduced to a couple of monolayers. [1, 2] Most of the properties of QPs rely
on their height so, in this work, we use a strain balance technique to further increase their aspect ratio
while keeping good crystal quality. [3]
The growth begins with a QDs layer formed after the deposition of 2 ML of InAs at 510 ºC on
a GaAs (001) substrate. On top of it, we grow a short period InAs/GaAsP superlattice at the same
substrate temperature. The GaAsP is 8.5 Å thick per period and has a nominal P content of 14 %. We
use 0.62 ML of InAs per period for sample A and 0.72 ML for sample B. The total number of periods
is 100 in both cases.
Figure 1 shows a detail of a transmission electron microscopy (TEM) image of sample A,
where it is clearly visible the formation of QPs. Their height is 120 nm and their diameter is around 13
nm, giving an aspect ratio of 9.2. Light emitted along the sample cleaved edge is linearly polarized
parallel to the growth direction, as expected from the large aspect ratio of the fabricated QPs (Fig 1b).
More surprisingly, light emitted along the growth direction is also strongly linearly polarized, as
shown in Fig 2c. This suggests that the short period P-based superlattice might be suffering from
composition modulation in the growth plane thus producing large in-plane strain anisotropy. On the
other hand, we see that increasing the In content of the QPs not only shifts the photoluminescence
emission to longer wavelengths, as expected, but also changes the relaxation dynamics from a pure
monoexponential decay at low In content to a faster stretched-like decay for higher In content (Fig. 2).
We suggest the appearance of very long living states or traps as responsible of this behavior.
Fig. 1: (a) TEM detail of a QP in sample A. Linear
polarization of light emitted along the (b) cleaved edge and
(c) growth direction.
Fig. 2: (a) PL of QPs with different In content and (b) their
corresponding time resolved PL decay curves.
[1] L. H. Li, P. Ridha, G. Patriarche, N. Chauvin, and A. Fiore, Appl. Phys. Lett, 92, 112102 (2008).
[2] H. J. Krenner, C. E. Pryor, J. He, and P. M. Petroff, Nanoletters, 8(6) 1750-1755 (2008).
[3] D. Alonso-Álvarez, B. Alén, J. M. Ripalda, A. G. Taboada, J. M. Llorens, Y. González, L. González, F. Briones, E.
Antolín, I. Ramirez, A. Martí, A. Luque, M. A. Roldán, J. Hernandez-Saz, M. Herrera and S.I.Molina, Proceedings 35th
IEEE Photovoltaic Specialist Conference, Hawaii (2010).
__________________________
* Contact: diego.alonso@imm.cnm.csic.es

We3.6
Polarity of GaN nanowires grown by Plasma-Assisted Molecular
Beam Epitaxy
K. Hestroffer 1* , C. Bougerol 1 , C. Leclere 2 , H. Renevier 2 , J. L. Rouvière 3 and B. Daudin 1
1 CEA-CNRS group « Nanophysique et Semiconducteurs », Institut Néel/CNRS-Univ. J. Fourier and CEA
Grenoble, INAC, SP2M, 17 rue des Martyrs, 38 054 Grenoble, France
2 Laboratoire des Matériaux et du Génie Physique, Grenoble INP - MINATEC, 3 parvis L. Néel 38016 Grenoble
3 CEA-Grenoble INAC/SP2M/LEMMA, 17 rue des Martyrs, 38054 Grenoble Cedex 9, France
GaN nanowires and related nanowire heterostructures are particularly attractive for
optoelectronic applications due to their defect free characteristics and the large range of wavelengths
covered by the different III-N alloys. As all the III-nitride structures crystallizing in the wurtzite phase,
the optical properties of these structures are governed by the presence of large internal electric fields.
On that account, the optical properties are crucially determined by the polarity of the structure, i. e.
whether the III-N bound is aligned in the growth direction or opposite to this one. In particular, as
regards the optical properties of thin GaN nanodisks inserted in AlN NW sections [1].
If the polarity of GaN 2D layers has already been widely investigated and understood [2, 3], it
is not the case of GaN nanowires. As a matter of fact, the characterization tools used to investigate the
2D layers lead to ambiguous results in the case of nanowires structures grown by MBE regarding their
small size. Besides, the polarity issue raises the question of the nanowires nucleation process and its
dependence on the substrate material or on a potential buffer layer.
Using a combination of four different characterization methods, namely Convergent Beam
Electron Diffraction, KOH chemical etching, direct observation by High Resolution Scanning
Transmission Electron Microscopy (HR-STEM) and anomalous X-Ray diffraction performed at the
European Synchrotron Radiation Facility, we demonstrate that GaN nanowires grown by PAMBE on
Si (111) are N-polar. We furthermore show that the presence of an AlN buffer layer does not affect the
polarity of the nanowires. On that account, we explain these observations by a possible nitridation of
the silicon surface prior to the nanowires nucleation and confront the results to the case of nanowires
grown on c-plane sapphire substrates.
Fig. 1: SEM image of GaN nanowires grown by PAMBE on bare Si (111) before (a) and after KOH etching (b).
HR-STEM image of a GaN nanowire (c) compared to the scheme of a GaN crystal projected on the (11-20) plane (d) (taken
from [5]) to identify the N-polarity.
[1] D. Camacho Mojica and Y. M. Niquet, Phys. Rev. B, 81 195313 (2010)
[2] B. Daudin et al., Appl. Phys. Lett. 69, 2480 (1996)
[3] A. R. Smith et al., Appl. Phys. Lett. 72, 2114 (1998)
[4] R. Armitage and K. Tsubaki, Nanotechnology, 21, 195202 (2010)
[5] J. L. Rouvière et al., Appl. Phys. Lett. 92, 201904 (2008)
*Contact: karine.hestroffer@cea.fr

2010
Varian becomes a part of Agilent
Varian Vacuum Technologies becomes part of Agilent Technologies
Two of Silicon Valley’s champions have combined to create a one-stop, truly global vacuum
supplier with a complete range of products and services for scientific research and industrial
applications.
• Varian’s 60+ years of technology leadership and tradition of innovation in vacuum and leak
detection technology become a part of Agilent, the world’s premier measurement company.
• Our shared values: customer satisfaction, commitment to quality, ethical standards, and
technology leadership.
• You can continue to rely on the same people, products, and services to assist you in solving
your vacuum challenges.
• As a leader in high- and ultra-high vacuum solutions for science and industry, Agilent is your
vacuum resource.
Visit us at www.agilent.com/go/varian.
Canada and US Toll Free Number: 800.882.7426 Europe Toll Free Number: 00.800.234.234.00

Monday Poster Session

MoP01
Post-heat treatment on the improvement of efficiency in CdTe
thin film solar cells
Zhizhong Bai, Deliang Wang*
Hefei National Laboratory for Physical Sciences at Microscale, University of Science and Technology
of China, Hefei, Anhui 230026, People’s Republic of China
High quality CdS and CdTe thin films were prepared by chemical bath deposition and
close-spaced sublimation, respectively. Post-heat treatment with the existence of
CdCl 2 is considered to be a ctitical process for obtaining a high efficiency CdTe/CdS
solar cell. In order to study the effect of post-heat treatment on CdTe films and
CdTe/CdS solar cells, three techniques were designed and investigated: CdTe/CdS
devices post-heat treated with a dip-coated CdCl 2 layer, devices post-heat treated with
presence of CdCl 2 powders, and devices post-heat treated with a vacuum evaporated
CdCl 2 coating layer. Vacuum evaporated CdCl 2 coating layer showed a uniform
morphology and provided a high vapor pressure of CdCl 2 during the heat treatment.
High CdCl 2 pressure enhanced p-type doping of the CdTe layer and a lowering of the
series resistance of the device. It also promoted CdTe recrystallization, grain growth,
and interdiffusion at the CdTe/CdS interface. Device with this post-heat treatment
process shows lager fill factor and higher efficiency. The backcontact of CdTe/CdS
devices was investigated mainly by changing the copper thickness in the backcontact
Cu/Au bilayer. Different copper thickness resulted in different buffer layers between
the CdTe film and the backcontact layer. An optimized CdCl 2 heat treatment and
backcontact layer deposition improved devices performance significantly. A
CdTe/CdS thin film solar cell with an AM1.5 efficiency of 13.2% were achieved.
Through the peak shifts of both Raman and XRD spectra, quantitative analysis on the
interdiffusion and its related reactions at the CdS/CdTe interface showed that ~ 11%
Te diffused into CdS and ~ 9% sulphur diffused into CdTe to form S-rich and Te-rich
CdS x Te 1-x alloy at the interface, respectively.
Fig.1(a)Cross-sectional SEM image of a CdTe/CdS device; (b) Current-voltage characteristic of the CdTe/CdS
solar cell with an efficiency of 13.2% under the standard AM1.5 condition.
______________________________________________________________________
* Contact: eedewang@ustc.edu.cn; Phone: +86-551-3600450

MoP02
Sn doped GaAs by CBE using tetramethyltin
C. García Núñez, D. Ghita, B. J. García *
Laboratorio de Electrónica y Semiconductores, Departamento de Física Aplicada, Universidad Autónoma de
Madrid, 28049 Madrid, Spain
Group-IV elements, when used as dopants, should be incorporated into the crystal lattice in
substitutional positions. In the case of GaAs:Sn, depending on the sublattice position Sn incorporates
to, it can behave either as a donor or an acceptor (amphoteric character), although Sn presents a strong
tendency to incorporate into the group-III sublattice, as a donor, under general epitaxy conditions. Sn
doped layers have been useful for GaInAs(P)/InP heterojunction bipolar transistors (HBTs), reducing
contact resistance in the emitter and collector regions, or within GaAs/AlGaAs heterojunction lasers.
High n-type doping levels can improve the operation frequency of these devices.
The highest carrier concentration level achieved for a given dopant is restricted by its solubility limit.
As an example, when Si is used as n-type dopant in GaAs, the above limit is n≈6x10 18 cm -3 [1]. The
use of Sn can increase that limit by near an order of magnitude (n≈1.5x10 19 cm -3 ) [2]. Previous
chemical beam epitaxy (CBE) works [3] used TESn as a Sn precursor, obtaining carrier concentrations
as high as n≈1x10 19 cm -3 . This work presents the results obtained on the CBE growth of GaAs:Sn,
using triethylgallium (TEGa), tertiarybutylarsine (TBAs), and tetramethyltin (TMSn), on GaAs(100)
substrates by CBE.
The samples were grown in a Riber 2300 MBE system modified to allow the use of gas sources [4].
Growth conditions were: growth rate v g =0.5 ml/s, substrate temperature T s =530 ºC, residual chamber
pressure 10 -6 mbar. Surface reconstruction and v g were determined by reflection high-energy electron
diffraction (RHEED). Considering similar sticking coefficients for all the atoms involved in the
growth process, the dopant flux should be 10 -4 -10 -7 lower than the fluxes of the main components. The
application of this requirement to our pressure regulated flux control system, will force it to work out
of range. We propose in this work the dilution of the Sn precursor in an inert carrier gas, such as H 2
(99.999% purity). Several series of samples with different dilution rates and dopant flux were grown.
The resistivity and type of carriers, their concentrations and mobilities (µ n =µ H /r H ) were determined by
Hall Effect measurements using the four-probe Van der Pauw technique. The Hall data have been
corrected for the effect of the surface depletion layer, due to the Fermi level pinning at the GaAs
surface originated by the high density of surface states; the thicknesses of the grown samples were
about 1 µm to reduce the influence of this effect. Figure 1 represents the obtained results on the n.
Four different GaAs:Sn sample series are plotted in this figure, with different dilution rates. All the
samples showed n-type behaviour, indicating that the residual p-type doping of our CBE grown GaAs
(n≈2x10 16 cm -3 due to C contamination) was compensated and even overcome by the Sn doping. As it
can be observed in the figure, n increases with both, the gas flux of the TMSn+H 2 mixtures, and the
TMSn content in the flux. The value n=1.6x10 19 cm -3 has been the highest doping obtained following
the described procedure. Therefore, as we can see in Figure 2, as n increases, µ n decreases. This
suggests that µ n in GaAs:Sn films is mainly dominated by the ionized impurity scattering (r H ≈1.9).
In summary, the electrical properties of the GaAs:Sn layers grown by CBE using H 2 -diluted TMSn
were investigated as a function of the TMSn dilution and the gas flux of the mixture. We showed that n
increases with TMSn concentration and gas flux, while µ n decreases with increasing n due to the
impurity scattering, as it is usually observed. Carrier concentrations up to n=1.6x10 -19 cm -3 have been
obtained, close to the Sn solubility limit, showing that the proposed doping method offers a precise
control of the n-doping in GaAs layers.
[1] C. R. Abernathy, S. J. Pearton, and N. T. Ha, J. Cryst. Growth 108, 827 (1991).
[2] C. R. Abernathy, S. J. Pearton, F. Ren, and J. Song, J. Cryst. Growth 113, 412 (1991).
[3] M. Weyers, J. Musolf, D. Marx, A. Kohl, and P. Balk, J. Cryst. Growth 105, 383 (1990).
[4] M. Aitlhouss, J. L. Castano, and J. Piqueras, Mat. Sci. Eng. B-Solid State Mat. Adv. Techn. 28, 155 (1994).
__________________________
* Contact: basilio.javier.garcia@uam.es

MoP02
(a)
(b)
Fig.1: Hall measurements from (a) carrier concentration, and (b) mobility, as a function of TMSn content in the gas,
and mixture flow rate.

MoP03
Early stages of the growth of InP and GaAs islands on
SrTiO3 substrates
B. Gobaut *1 , J. Penuelas 1 , A. Chettaoui 1 , J. Cheng 1 , G. Grenet 1 , L. Largeau 2 and
G. Saint-Girons 1
1
Institut des Nanotechnologies de Lyon, Ecole Centrale de Lyon, 36 av. G. de Collongue, 69134 Ecully, France
2 Laboratoire de Photonique et Nanostructure, Route de Nozay, 91460 Marcoussis, France
The integration of III-V semiconductor on silicon is of great interest for the development of integrated
optoelectronic properties on Si.
However, the direct epitaxy of III-V on Si is made difficult by the large lattice mismatch between the
two semiconductors which leads to a large density of defects in the epitaxial layer, especially threading
dislocations due to the plastic relaxation. Recently, we have demonstrated the great advantages of
using an interfacial layer of oxide grown on Si, typically SrTiO 3 , as a buffer layer. The lattice
mismatch between the III-V and the oxide layers is accommodated by a network of dislocations
confined at the heterointerface which allows the top layer to grow fully relaxed [1]. This so-called
compliant interface reduces drastically the density of defects in the epitaxial layer and thus improves
the properties of the device [2]. III-V semiconductors on STO are deposited according a Volmer-Weber
growth mode followed by a coalescence part. To further reduce the number of defects in the III-V
layer, we give a particular attention to the shapes and structures of the islands during the early stages
of the growth.
In this contribution, we will present our study of the early stages of the growth of InP and GaAs
islands deposited on STO/Si templates. Our results highlight the influence of the surface and growth
conditions on the crystalline orientation of III-V nano-islands. Another part of our study will also
include a description of the shapes and facets of the dots observed by Transmission Electron
Microscopy (TEM) and X-Ray diffraction and the crystalline form (cubic or wurtzite). Experiment of
in-situ monitoring of growth of Ge islands on STO by Grazing Incidence Small Angle X-ray
Scattering (GISAXS) performed at the synchrotron ESRF will also be presented.
(a)
Fig 1: Images of InP/STO/Si islands observed by STEM-HAADF (a) and by HRTEM (b)
(b)
[1] G. Saint-Girons, J. Cheng, P. Regreny, L. Largeau, G. Patriarche and G. Hollinger Physical Review B 80, 155308 (2009)
[2] B. Gobaut, J. Penuelas, J. Cheng, A. Chettaoui, L. Largeau, G. Hollinger and G. Saint-Girons, (submitted).
__________________________
* Contact: benoit.gobaut@ec-lyon.fr

MoP04
Growth directions and structural properties of InP nanowires
fabricated on Si and SrTiO 3 substrates
J. Penuelas 1 , K. Naji 1 , H. Dumont 1 , G. Saint-Girons 1 , G. Patriarche 2 , M. Gendry 1
1
Université de Lyon, Institut des Nanotechnologies de Lyon-INL, UMR CNRS 5270, Ecole Centrale de Lyon,
36 avenue Guy de Collongue, 69134, Ecully, France
2 Laboratoire de Photonique et de Nanostructures-LPN, UPR CNRS 20, Route de Nozay, 91460, Marcoussis,
France
Semiconducting nanowires (NWs) have attracted a lot of interest because of their application in
nanoelectronics [1], photonics [2], energy conversion [3] and biology [4]. NWs approach is also an
interesting way for monolithic integration of mismatched III-V semiconductors on silicon substrate.
Molecular Beam Epitaxy (MBE) using the Vapor-Liquid-Solid (VLS) method can provide control of
the NWs structure at atomic scale and also opens the way to grow complex one-dimensional
heterostructures (core-shell NWs, stacking of different semiconductors, etc…). However, a key point
for the NW monolithic integration is the control of the growth direction of the NWs and to be able to
produce NWs standing vertically on their substrate.
In this context, we have grown InP NWs on Si(001) and Si(111) substrates by VLS-MBE using gold
(Au) or gold-indium alloy (Au-In) as catalyst. We have also grown InP NWs on SrTiO 3 (001) substrate,
this oxide which can be an appropriate template on Si(001) to promote vertically standing InP NWs.
In this communication, we present results concerning the NWs structural properties studied by X-ray
diffraction and by Transmission Electron Microscopy. We report the observation of different
morphologies depending on the catalyst droplet diameter and substrate type. The NW structure also
strongly depends on the substrate temperature and V/III beam ratio. We evidence and discuss various
NW to substrate epitaxial relationships and a particular attention is given to explain the NWs growth
direction in conjunction with the epitaxial relationship. We demonstrate that X-ray diffraction is an
accurate tool to measure this growth direction.
(a)
(b)
χ= 90°
a 1
a 1
b
a 2
b
a 2 b
a 2
a2
b
χ = 30°
χ= 60°
a 1
φ
Si [010]
a 1
Si [001]
Si [100]
Fig 1 : InP NWs on Si(001) substrate: (a) Scanning Electron Microscopy image, (b) X-ray diffraction pole figure around
(0004) reflection of InP (the spots a1 and a2 are related to InP NWs having different growth directions, b spots are related to
Si).
[1] Jean-Pierre Colinge, Chi-Woo Lee, Aryan Afzalian, et al., Nature Nanotechnology 5 225 (2010).
[2] Ruoxue Yan, Daniel Gargas, Peidong Yang, Nature Photonics, 3, 569 (2009).
[3] Guang Zhu, Rusen Yang, Sihong Wang, and Zhong Lin Wang, Nano Letters, 10, 3151 (2010).
[4] Alex K. Shalek, Jacob T. Robinson, Ethan S. Karp et al. PNAS, 107, 4489 (2010)
________________________
* Contact: jose.penuelas@ec-lyon.fr

MoP05
Optimization of MBE-grown AlSb/InAs High Electron Mobility
Transistor Structures
H. Zhao 1* , G. Moschetti 2 , S. Wang 3 , P-Å. Nilsson 2 , and J. Grahn 2
1
Terahertz and Millimetre Wave Laboratory, 2 Microwave Electronics Laboratory, 3 Photonics Laboratory,
Department of Microtechnology and Nanoscience, Chalmers University of Technology, SE-41296, Göteborg,
Sweden
The AlSb/InAs high electron mobility transistor (HEMT) is a promising device candidate for low
power consumption and high speed applications. However, due to the larger lattice constants of AlSb
and InAs with respect to commercial InP and GaAs substrates, a thick and relaxed AlSb metamorphic
buffer is necessary to act as a template. In this work, AlSb/InAs HEMT structures are grown on both
GaAs and InP substrates. Surface morphology of the AlSb metamorphic buffer depends on the type of
substrate and the growth temperature. The optimized AlSb buffer on InP substrate shows anisotropic
and larger surface roughness compared with that grown on GaAs, as shown by AFM in Fig.1. Since
the AlSb/GaAs has a larger thermal expansion coefficient difference and a larger lattice mismatch
compared with the AlSb/InP, the surface morphology is most likely related to the initial strain relaxation
near the AlSb/substrate interface and to the different diffusion lengths of adatoms along the orthogonal
crystallographic directions [1]. Furthermore, the surface roughness of the AlSb buffer has an important
effect on the AlSb/InAs channel interfaces, thus affecting the electron mobility.
The electron mobility of bulk InAs grown on an AlSb metamorphic buffer was first investigated at a
growth temperature range of 450-530 o C. Higher electron mobility was observed at a higher As
cracking temperature, indicating the growth of InAs prefers to the As 2 mode instead of the As 4 mode.
The electron mobility of InAs is more than twice higher when grown at 510 o C than that grown at
temperatures below 500 o C, but also decreases rapidly when the growth temperature increases to 530
o C. Therefore the growth temperature window for InAs is quite narrow.
Finally, the AlSb/InAs HEMT structure, see Fig.2, has been optimized on GaAs substrates. An
optimal AlSb spacer layer of 6 nm was found resulting in a maximum electron mobility of 1.9x10 4
cm 2 /Vs as shown in Fig.3.
[1] Li, H., Q. Zhuang, Z. Wang, and T. Daniels-Race, J. Appl. Phys., 87, 188 (2000).
_________________________
* Contact: zhaoh@chalmers.se

MoP05
Fig.1: AFM measurements performed on top of the AlSb buffer grown on InP (left) and GaAs substrates (right). It shows
RMS roughness of 0.732 nm and 1.596 nm on InP and of 0.449 nm and 0.674 nm on GaAs, along the [1-10] and [110]
directions respectively.
Fig.2: Typical HEMT structure with Si δ-doping (dash line).
Fig.3: The electron mobility and sheet density of HEMT grown on GaAs as a function of the AlSb spacer layer thickness.

MoP06
Monolithic integration of InP based heterostructures on silicon using
SrTiO 3 templates
A. Chettaoui 1* , B. Gobaut 1 , J. Penuelas 1 , J. Cheng 1 , G. Niu 1 , L. Largeau 2 , P. Regreny 1 , G.
Saint-Girons 1
1 Université de Lyon, Institut des Nanotechnologies de Lyon (UMR5270/CNRS), Ecole Centrale de Lyon, 36
avenue Guy de Collongue, 69134 Ecully cedex, France
2 LPN/UPR20-CNRS, Route de Nozay, 91460 Marcoussis (France)
The integration of various materials on Si substrate is a key issue for the future development of
complementary metal oxide semiconductor (CMOS) technologies. In particular, integrating Ge and
III-V on Si could allow combining optoelectronic functionalities with standard Si-based CMOS
systems.
In this context, we have recently demonstrated and studied the compliant behaviour of the III-V/
SrTiO 3 interfaces [1,2]. The large mismatch between III-V and SrTiO 3 is fully accommodated by a
network of geometric dislocations confined at the heterointerface. Consequently, the growing material
takes its bulk lattice parameter as soon as growth begins, does not undergo any plastic relaxation
mechanism and is free of extended defects related to plastic relaxation.
In this contribution, we will show how this specific accommodation pathway can be used to grow
InAsP/InP heterostructures on SrTiO 3 /Si templates. In particular, we will present a detailed study of
the influence of the growth parameters of InP on SrTiO3 on the structural quality of the III-V material
[3]. Moreover; we will present a comparative study of the optical and structural properties of
InAsP/InP quantum well heterostructures grown on SrTiO 3 /Si templates, or directly on Si without
oxide template [4,5].
Finally, we will propose some perspectives and discuss the potential of this approach for III-V
semiconductors integration on Si.
This work is supported by the “Agence Nationale de la Recherche” (project P3N2009 « COMPHETI »
#ANR-09-NANO01301) and by the Rhône-Alpes region (projet cluster micronano « MonoSi »).
* Contact: azza.chettaoui@ec-lyon.fr

MoP06
2/ (°)
Normalized intensity (a.u.)
71
69
67
65
63
61
59
InP 331
(a)
InP 004
Si 004
Si
004
InP
004
(b)
-10 -8 -6 -4 -2 0 2 -2 0 2
1.0
0.8
0.6
0.4
0.2
(c)
(°) (°)
Sample A
Sample B
0.0
-2 -1 0 1 2
(°)
Fig. 1 (a) X-Ray diffraction reciprocal space map recorded on an
InAsP/InP heterostructure grown on Si (sample B). (b)
Reciprocal space map recorded on the same heterostructure
grown on a SrTiO 3 /Si template (sample A); (c) ω-scans of the InP
004 peak of sample A (red cut line in (b)) and sample B (blue cut
line in (a)).
Intensity (a.u.)
InP
STO
Si(001)
0.8
300 K
0.6
0.4
0.2
0.0
Sample B
(x50)
Sample A
(a)
(b)
1400 1500 1600 1700
Lambda (nm)
Fig. 3 (a) TEM cross-sectional view of sample A. The
quantum well is indicated by an arrow. (b)
Photoluminescence measurements for sample A and B.
[1] G. Saint-Girons, J. Cheng, P. Regreny, L. Largeau, G. Patriarche and G. Hollinger, Phys. Rev. B 80, 155308, (2009).
[2] G. Saint-Girons, C. Priester, P. Regreny, G. Patriarche, L. Largeau, V. Favre-Nicolin, G. Xu, Y. Robach, M.
Gendry ; G. Hollinger Appl. Phys. Lett. 92, 241907 (2008).
[3] J. Cheng, L. Largeau, G. Patriarche, P. Regreny, G. Hollinger, G. Saint-Girons, Appl. Phys. Lett. 94, 231902, (2009)
[4] J. Cheng, T. Aviles, A. El-Akra, C. Bru-Chevallier, L. Largeau, G. Patriarche, P. Regreny, A. Benamrouche, Y. Robach,
G. Hollinger and G. Saint-Girons, Appl. Phys. Lett. 95, 232116, (2009).
[5]B. Gobaut et al, to be published in applied physics letters.
* Contact: azza.chettaoui@ec-lyon.fr

MoP07
Crystal structure X-ray investigation of InAs nanorods on Si(111)
A. Davydok 1* , A. Biermanns 1 , M. Dimakis 2 , S. Breuer 2 , L. Geelhaar 2 , and U. Pietsch 1
1
Festkörperphysik, Universität Siegen, Walter-Flex-Str. 3,57072, Siegen, Germany
2 Paul-Drude-Institut für Festkörperelektronik, Hausvogteiplatz 5-7,10117 Berlin, Germany
Semiconductor nanowires (NWs) are of particular interest for new semiconductor devices.
One future application is the development of one-dimensional field effect transistors (FET) with small
quantum capacity, improved power scaling and ideal linearity. For application as light emitting diodes
(LED) it is promising that the nanowire approach can be used to form heterostructures of materials
with a large lattice mismatch.
Group III-V nanowires offer the exciting possibility of epitaxial growth on a wide variety of
substrates, most importantly silicon. One of the most common methods for NW growth is the Vapour-
Liquid-Solid(VLS) mechanism. It was found that by MBE-mode of VLS method nearly any A III B V
semiconductor material can be grown as NWs onto another A III B V or group IV [111] substrates
independent from lattice mismatch.
Using x-ray diffraction we have investigated InAs NW’s on Si[111] substrate where the lattice
mismatch is about 11%. Considering our experience in GaAs NW’s structure analysis [1,2] we have
performed X-ray inspection of InAs nanowires on different stages of growth process using
synchrotron radiation provided by ESRF. NW’s ensemble measurements have been done at selected
Bragg reflections. Using symmetrical and asymmetrical Bragg reflections which are particularly
sensitive to wurzite and/or zinc-blende structure we were able to extract information about the strain
accommodation between substrate and NW. With help of the nano-focus setup of ID1 beamline at
ESRF we were able to characterize single NW´s grown at one and same sample. Figure 1(a) shows
reciprocal space maps of 3 samples grown by MBE where the growth was stopped after 5 sec
(nanowires length is ~35nm), 10 sec (nanowires length is ~75nm) and 43 sec (nanowires length is
~250nm). Using the nano-focus setup at ID1 we have performed measurements of selected single
nanowires (Fig.1b). It is obvious that the peak separation between the Si substrate and the InAs NW
changes as function of growth time. Peak position slight shift during growth process can be explained
by WZ-ZB transition. By dashed line InAs (111) position is marked.
(a)
Fig 1: (a) Reciprocal space map at (111) reflection;
(b) Examples of single nanowire measurements of 3 samples at (111) reflection
(b)
[1] A. Biermanns, A. Davydok, H. Paetzelt, A. Diaz, V. Gottschalch, T.H. Metzger and U. Pietsch, J. Synchrotron Rad.
(2009). 16, 796-802
[2] A. Davydok, A. Biermanns, U. Pietsch, J. Grenzer, H. Paetzelt and V.Gottschalch Metallurgical and Materials
Transactions A: 41, 5, 1191 (2010)
__________________________
* Contact: davydok@physik.uni-siegen.de

MoP08
Structural and optical properties of InN films grown on ZnO(0001)
by plasma-assisted molecular beam epitaxy
Y. J. Cho * , O. Brandt, M. Ramsteiner, M. Wienold and H. Riechert
Paul-Drude-Institut für Festkörperelektronik, Hausvogteiplatz 5–7, 10117 Berlin, Germany
The narrow band gap (~0.65 eV) semiconductor InN is the least understood of the group III
nitrides. Fundamental studies are often hampered by the fact that the crystal quality of InN lags behind
other group III nitrides primarily because of the very low decomposition temperature of InN.
The group II-VI semiconductor ZnO, being of identical crystal structure as the group-III nitrides,
is an interesting substrate for the growth of InN, as it is available as bulk crystal with high quality and
excellent surface finish even for its (000-1) (O) face. Indeed, it has recently been reported that
high-quality InN films with a full-width-at-half-maximum (FWHM) of the (002) x-ray rocking curve
(XRC) of 150″ can be grown on ZnO(000-1) substrates [1]. However, a chemical reaction between the
two materials is activated with increasing substrate temperature T s and results in intermixed interface
layers as well as in a severe out-diffusion of O, which is acting as donor in overgrown InN layers.
Figure 1(a) shows the FWHM of x-ray (002) and (102) XRCs and RHEED patterns of ~130 nm
thick InN films grown on ZnO(000-1) at different T s . High T s is seen to markedly improve the
morphology and simultaneously results in a narrowing of the symmetric (002) XRC, eventually leading
to a value of only 27″ for T s = 450 °C. While the density of screw dislocations is thus very low for these
layers, that of edge dislocations does not show any noticeable change with T s and remains very high.
Figure 1(b) depicts photoluminescence (PL) spectra recorded at 300 K for these films. Contrary to the
evolution of the morphology and structural quality with T s , a monotonic decrease of PL intensity with
increasing T s is observed. The room-temperature PL is found to be completely quenched for samples
grown above 450 °C.
FWHM 002 (arcsec)
400
300
200
100
0
350 o C
350 o C
450 o C
550 o C
350 o C
350 400 450 500 550
T s ( o C)
3000
2800
2600
2400
2200
FWHM 102 (arcsec)
(a)
Intensity (arb. units)
300 K
350 o C
400 o C
450 o C
0.5 0.6 0.7 0.8 0.9 1.0
E (eV)
Fig 1: (a) FWHM of the (002) and (102) reflections in XRCs, and (b) room temperature PL spectra of ~130 nm thick
InN on ZnO(000-1) as a function of substrate temperature. The inset in (a) shows the RHEED patterns of the InN
films grown at the temperatures indicated.
Figure 2(a) shows the evolution of the FWHM of (002) and (102) XRCs on the thickness of the
InN films grown at T s = 425 °C under slightly N rich condition. A dramatic reduction of the FWHM of
the (102) XRCs with increasing InN thickness is observed, implying an efficient annihilation of edge
dislocations. The density of screw dislocations seems not to depend on thickness. Moreover, the PL
intensity increases and the PL peak energy decreases with increasing film thickness as seen in Fig. 2(b).
The 2 µm thick InN film exhibits very satisfactory value for both screw and edge dislocations, a decent
morphology [see the inset of Fig. 2(a)] and strong room-temperature PL at the band gap of InN. Guided
by these results, we develop a growth strategy to obtain high-quality InN films by PAMBE with smooth
surfaces, high structural quality and an PL efficiency comparable to that of GaN.
[1] T. Ohgaki et al., J. Cryst. Growth, 292, 33 (2006).
_________________________
* Contact: yjcho@pdi-berlin.de
(b)

MoP08
FWHM 002
(arcsec)
500
400
300
200
100
0
0 500 1000 1500 2000
t (nm)
(a)
2200
2000
1800
1600
1400
1200
1000
800
600
FWHM 102
(arcsec)
Intensity (arb. units)
300 K
2000 nm
500 nm
140 nm
0.5 0.6 0.7 0.8 0.9
E (eV)
Fig 2: InN thickness dependent (a) FWHM of the x-ray rocking curves for (002) and (102) reflections, and (b) room
temperature PL spectra. Inset in (a) shows the RHEED patterns of the 2000 nm thick InN film after growth.
(b)

MoP09
Multifunctional Epitaxial Nanocomposite Films by LMBE
J.Xiong 1,*
1
Key Lab of Electronic Thin Films and Integrated Devices, University of Electronic Science and Technology of
China, Chengdu 610054, People’s Republic of China
Transition metal oxides have attracted great attention due to their versatile properties. Various
composite, multilayers, and artificial superlattices have been demonstrated different excellent multifunctionality
or enhanced unique physical properties over single layer pure oxide thin films at
multifarious fields. The coupling and interactions in the nanocomposited and multilayered systems can
strongly alter the films’ growth features and their physical properties. We have made tremendous work
in the preparation and characterization of extremely high quality metal-oxide thin films using
molecular beam epitaxy (MBE) technique. Specifically, we have successfully prepared
[(BiFeO 3 ) n /(La 0.7 Sr 0.3 MnO 3 ) n ] m and [(BiFeO 3 ) n /(BiMnO 3 ) n ]m superlattices and single-crystal strained
nanocomposite BiFeO 3 (BFO) 0.5 : BiMnO 3 (BMO) 0.5 films, and observed that there is a strongly
improvement in magnetic properties compared to the pure BFO films. We investigated and compared
three particular epitaxial nanocomposite architectures: the vertically aligned nanocomposite, the
layered laminar-like nanocomposite, and the mixed nanocomposite, and understand how the strain
affects the properties of ferroelectric, dielectric, magnetic and optical. We plan to explore various new
phenomena in the as-desired nanocomposited or multilayered structures and to pave a way for the new
concept device development.
__________________________
* Contact: jiexiong@uestc.edu.cn

MoP10
Near infrared high efficiency InAs/GaAsSb QDLEDs: band
alignment and carrier recombination mechanisms
A. Hierro * , M. Montes, M. Moral, J.M. Ulloa, A. Guzman
ISOM and Dpto. Ing. Electronica, Univ. Politecnica Madrid, Ciudad Universitaria s/n, 28040 Spain
The development of high efficiency laser diodes (LD) and light emitting diodes (LED) covering the
1.0 to 1.55 µm region of the spectra using GaAs heteroepitaxy has been long pursued. Due to the lack
of materials that can be grown lattice-macthed to GaAs with bandgaps in the 1.0 to 1.55 µm region,
quantum wells (QW) or quantum dots (QD) need be used. The most successful approach with QWs
has been to use InGaAs, but one needs to add another element, such as N, to be able to reach
1.3/1.5µm. Even though LDs have been successfully demonstrated with the QW approach, using N
leads to problems with compositional homogeneity across the wafer, and limited efficiency due to
strong non-radiative recombination. The alternative approach of using InAs QDs is an attractive
option, but once again, to reach the longest wavelengths one needs very large QDs and control over
the size distribution and band alignment. In this work we demonstrate InAs/GaAsSb QDLEDs with
high efficiencies, emitting from 1.1 to 1.52 µm, and we analyze the band alignment and carrier loss
mechanisms that result from the presence of Sb in the capping layer.
The devices consist of a p-i-n structure where a single layer of InAs/GaAsSb QDs is immersed in the
intrinsic region. The QDs were always formed by depositing 2.7 MLs of InAs under the same growth
conditions, and the only variation was the Sb content in the 4 nm thick capping layer, which was
changed from 0 to 28 %, as measured by cross-sectional scanning tunneling microcopy (X-STM) [1].
The p and n regions were doped with 2x10 18 cm -3 Be and Si, respectively. Mesa structures were
defined by wet chemical etching down to the buffer layer to electrically isolate the devices. Top ohmic
ring contacts provided proper light extraction through the front surface, whereas a back ohmic contact
was deposited on the entire substrate. The devices were mounted on TOs and wire bonded.
As shown in Fig. 1, as the Sb content in the capping is increased, the QDLEDs cover successfully the
1.1 to 1.52 µm region at room temperature. There is also an increase in the electroluminescence line
width for wavelengths above 1.4 µm that could arise from increased QD size inhomogeneity.
However, X-STM and AFM analysis show that this is not the case. Indeed, analysis of the
electroluminescence emission as a function of injected current shows very different behaviors below
and above 1.4 µm. Below 1.4 µm, both the ground state and up to two excited states can be observed,
with a clear saturation of the lowest energy states at high currents, and negligible blue shifts,
characteristic of a type I band alignment. Above 1.4 µm, at large currents, the ground state emission
blue shifts strongly with current (Fig. 2), characteristic of a type II band alignment. However, even
with a type II band alignment, the QDLEDs present an external efficiency close to the type I QDLEDs,
and actually much larger than that of the reference Sb-free QDLED emitting at 1.15 µm (Fig. 3).
The analysis of the electroluminescence thermal quenching shows for the type I QDLEDs an
activation energy for the ground state decreasing from 225 to 100 meV, for wavelengths shifting from
1.1 to 1.3 µm (Fig. 4). This decrease is consistent with the decreased valence band offset as the Sb is
increased, and indicates that hole leakage from the QDs to the capping layer is the dominant carrier
loss mechanism. In the case of the type II QDLEDs, the activation energy depends strongly on the
injected current. At low currents, where the bands have a large slope due to the junction built-in field,
the activation energy is quite large, around 275 meV, likely arising from hole leakage from the capping
layer to the GaAs barrier. However, at high currents, where flat band conditions are achieved, the
activation energy decreases down to ~150 meV (Fig. 4), which can be explained to arise from electron
leakage from the QD excited states to the GaAs barrier and/or GaAsSb capping. Taking together the
activation energies and the emission energies of the QDLEDs allows one to define the structure band
alignment as a function of capping Sb content and wavelength.
[1] J.M. Ulloa et al., Phys.Rev.B. 81,165305, 2010.
______________
* Contact: adrian.hierro@upm.es

MoP10
Normalized EL intensity
(arb. units)
0.64 A·cm -2 1.0
80
J=318 A·cm -2
0.8
60
Type I
0.6
40
0.4
0.2
20
Type II
0.0
0
1.0 1.1 1.2 1.3 1.4 1.5 1.6
1100 1200 1300 1400 1500
Wavelength (µm)
GS blue shift (meV)
GS wavelength (nm)
Fig 1: RT electroluminescence spectra from the QDLEDs.
Fig 2: Ground state blue shift as a function of the emission
wavelength at high injection currents.
External efficiency (mW/A)
0.35
0.30
Circles: maximum η ext
Triagles: η ext
at 1.6 A cm -2
0.25
0.20
0.15
0.10
Type I Type II
1100 1200 1300 1400 1500
GS transition wavelength (nm)
Fig 3: External efficiency of the QDLEDs at low injection
current for the ground state emission.
300
E act
(meV)
200
100
Type I
Type II
1100 1200 1300 1400 1500
Wavelength (nm)
Fig 4: Electroluminescence thermal activation energy from
the QDLEDs obtained between 200 and 300K under high
injection current conditions.

MoP11
In-situ Reflectance Anisotropy Spectroscopy (RAS) for doping
control during MBE growth of AlGaInAsSb laser structures
D. Hoffmann * , T. Loeber and H. Fouckhardt
Department of Physics, Integrated Optoelectronics and Microoptics Research Group, Kaiserslautern University
of Technology , P.O. Box 3049, D-67653 Kaiserslautern, Germany
In this work reflectance and reflectance anisotropy spectroscopy (RAS) [1-3] is employed as in-situ
monitoring technique to study the growth of AlGaInAsSb lasers. The RAS signal derives from a
reconstructed surface where an existing anisotropy along two principal axes leads to a difference
signal between the optical reflectance R along these axes. This difference signal can be measured as a
function of photon energy.
The combined measurement of reflectance and reflectance anisotropy spectra offers extensive growth
information during molecular beam epitaxy (MBE). For example the spectroscopic reflectance
contains information on the growth rate, the morphology and the composition of ternary layers. In
addition the RAS signal allows for monitoring oxide desorption, surface reconstruction during growth,
surface response to growth interruptions, and doping levels.
All devices are grown in a DCA R450 MBE with standard effusion cells for group-III-elements (Al,
Ga, In), and valved crackers for group-V-elements (Sb, As). Te and Be effusion cells are utilized for n-
and p-doping, respectively. RAS measurements are done with the EpiRAS ® TT-system by LayTec.
Two different types of diode lasers are grown. The first one is realized on n-GaSb substrates with n-
type and p-type AlGa(As)Sb claddings - lattice matched to GaSb - and an undoped active region with
a (Ga)InAs(Sb)/GaSb-multiple quantum well (MQW). The latter is surrounded by undoped GaSb
layers for waveguide broadening. The second type of diode lasers contains stacked GaAsSb-quantum
dot (QD) layers within GaAs barriers. Here the laser growth is based on GaAs-substrates and the
cladding layers consist of n- and p-type AlGaAs, respectively.
RAS is applied for monitoring the essential oxide desorption before the main growth process.
Afterwards the colorplot measurement mode (full spectrum taken, RAS signal color-coded) is used for
control of the entire growth process, leading to a fully spectroscopic fingerprint of the sample (shown
in Fig. 1). This way an unusual behavior during growth, for example in terms of an anomalous doping
of cladding layers, can be directly unveiled. Furthermore the fingerprint proves the reproducibility of
the individual wells in the MQW structure.
The qualitative influence of Be and Te as p- and n-dopant in the AlGaSb and AlGaAs cladding layers
on the RAS signal is revealed (shown in Fig. 2). In addition, the effect of doping is confirmed by exsitu
measurements such as SIMS and C/V for a quantitative understanding.

MoP11
Fig 1: RAS fingerprint of an entire growth of a laser sample in comparison to the schematic layer design. The
measurement signal is color-coded (blue = low signal, red = high signal) and plotted against photon energy (abscissa)
and time (ordinate). On the right the enlarged RAS colorplot of the active zone is shown for proof of reproducibility.
Fig 2: Doping effects in cladding layers: (a) RAS spectra of undoped and n-doped AlGaSb, (b) time resolved
measurement during growth of AlGaSb, undoped and with varying doping concentrations (source temperature of Be
is given for each layer as reference), (c) time resolved measurement during growth of AlGaAs, undoped, n- and p-
doped with varying doping concentrations (source temperatures of Te and Be are given as reference).
[1] V. L. Berkovits, JETP Lett., 41, 551 (1985).
[2] D. E. Aspnes and A. A. Studna, Phys. Rev. Lett., 54, 1956 (1985).
[3] D. E. Aspnes, J. P. Harbison, A. A. Studna, and L. T. Florez, J. Vac. Sci. Technol., A 6, 1327 (1988).
__________________________
* Contact: dhoffmann@physik.uni-kl.de

MoP12
Effect of growth temperature on surface morphology of
selectively grown GaN layers by ammonia-based metal-organic
molecular beam epitaxy
C. H. Lin, R. Abe, S. Uchiyama, Y. Uete, T. Maruyama and S. Naritsuka*
Department of Materials Science and Engineering, Meijo University,
1-501 Shiogamaguchi, Tenpaku-ku, Nagoya 468-8502, Japan
III-nitride materials had been investigated broadly because they have a very wide range of band-gap
energy which offers high potential for fabricating a variety of photoelectric devices. Superior physical
characteristics such as large bond strength also permit that III-nitride devices can be operated at high
temperature and with high supplied voltage. Further, no toxic sources are used during the production,
compared with other III-V compound materials. Nevertheless, the quality of III-nitrides still needs to be
improved since they have large numbers of threading dislocations, which deteriorate the performances
of the devices.
Microchannel epitaxy (MCE) is considered as an effective technique for reducing threading
dislocations, and dislocation-free regions are obtained in the laterally grown areas [1]. Therefore, we
have recently studied MCE of GaN using ammonia-based metal organic molecular beam epitaxy
(NH 3 -based MOMBE) [2-3]. The results show that the surface morphology of the grown layers, which
are largely changed by the growth temperature, strongly affects the length of the accompanying lateral
growth of GaN. Therefore, in this paper the effect of growth temperature on the surface morphology of
selective growth of GaN is systematically investigated. The mechanism of the lateral growth is also
studied in concerns with the inter-surface diffusion of the ad-atoms.
The surface morphologies of GaN selectively grown layers measured by atomic force microscopy
(AFM) are shown in Fig.1. The increase of the growth temperature is found to improve the surface
smoothness of the layer. The X-ray photoemission spectroscopy (XPS) measurements also indicate the
presence of Ga atoms on the surface of the layers grown at low growth temperatures. The cross-sectional
SEM image of the sample grown at 820°C shows that the selective growth was accompanied with the
lateral growth. This suggests that the inter-surface diffusion of the ad-atoms was occurred from the top
surface to the side surfaces, which was possibly produced not only by the high growth temperature but
also by the formation of the flat surface on the top of the growth.
650 o C 700 o C 820 o C 820 o C
Inter-surface diffusion
RMS: 31 nm RMS: 14 nm RMS: 6 nm
(a) AFM
(b) Cross-sectional SEM
Fig 1: Surface morphologies of GaN selective grown layers
3µm
[1] T. Nishinaga, T. Nakano and S. Zhang, Jpn. J. Appl. Phys. 27 L964 (1988).
[2] C. H. Lin, R. Abe, T. Maruyama and S. Naritsuka, “Temperature dependence of selective growth of GaN by ammonia-based
metal-organic molecular beam epitaxy”, J. Crystal Growth, in print.
[3] C. H. Lin, R. Abe, T. Maruyama and S. Naritsuka, “Low angle incidence microchannel epitaxy of GaN grown by
ammonia-based metal-organic molecular beam epitaxy”, J. Crystal Growth, in print.
* Contact: narit@meijo-u.ac.jp

MoP12
Normalized Intensity (a.u.)
1.0
0.8
0.6
0.4
0.2
o
C
o
C
o
C
820?
700?
650?
Ga-N
Ga-O
0.0
14 15 16 17 18 19 20 21 22 23
Binding Energy (eV)
Fig.A XPS spectra of selectively grown GaN layers
Fig.B Growth temperature dependence of intensity
ratio of Ga3d/N1s
(a)
(b)
(c)
(d)
Fig.C SEM images of selective growth of GaN grown at 650 o C (a), 700 o C (b), 750 o C(c) and 820 o C (d) with 5µm/5µm
(windows/masks) stripes. The inset in each figure shows the corresponding RHEED pattern (e-beam//[11-20]).

MoP13
Investigation of the local electronic structure of Cu-doped GaN
grown by plasma assisted MBE
R. Schuber 1,* , P.R. Ganz 1 , F. Wilhelm 2 , A. Rogalev 2 , and D.M. Schaadt 1
1
Institute of Applied Physics/DFG-Center for Functional Nanostructures (CFN), Karlsruhe Institute of
Technology (KIT), 76131 Karlsruhe, Germany
2 European Synchrotron Radiation Facility (ESRF), 38043 Grenoble Cedex, France
Dilute magnetic semiconductors (DMS) are a topic of great interest in today’s research. Many material
systems have been investigated with the aim to obtain a high temperature ferromagnetic
semiconductor. The material system of GaN doped with transitional elements has been theoretically
predicted to be ferromagnetic at room temperature [1,2]. GaN doped with the intrinsically
nonmagnetic element Cu has been reported to indeed exhibit ferromagnetic behavior at room
temperature, in implanted films [3], nanowires [4] and in MBE grown films [5]. However, there are
yet many unanswered questions concerning the mechanism of ferromagnetism in this system. Above
all, a detailed understanding of the incorporation of Cu in the GaN host is desirable.
A way of probing the local electronic structure of Cu-doped GaN is given by element specific x-ray
absorption spectroscopy (XAS). Especially x-ray linear dichroism (XLD) is a suitable technique to
explore the local structure and symmetry of a material by probing the anisotropy of the electronic
structure [6]. We therefore measured the XLD as well as the x-ray absorption near edge structure
(XANES) at the Cu and Ga K-edges of GaN:Cu wurtzite DMS samples at the ID12 beamline of the
European Synchrotron Radiation Facility (ESRF). In this way we investigated GaN:Cu samples with
nominal Cu concentrations between 0% and 2.3%.
The samples grown by plasma assisted MBE in Ga rich conditions show the formation of Cu 9 Ga 4
compounds (identified by transmission electron microscopy analysis) on their surfaces. To remove
these compounds from the surfaces the samples were etched with HNO 3 for 5 min. XAS spectra were
measured for etched and unetched samples.
For clarification of the experimental results, i.e. to clarify the role of the surface compounds and to
evaluate the Cu site position in the GaN host, i.e. Cu on Ga sites, N sites or interstitial sites, we
performed simulations of the GaN:Cu and the Cu 9 Ga 4 crystals for the Cu and Ga K-edges at different
doping levels using the FDMNES code [7]. A comparison with the experimental results shows that the
Cu atoms predominantly occupy Ga and interstitial sites.
[1] C. Liu et al., J. Mater Sci. - Mater: Electron., 16, 555 (2005).
[2] R.Q. Wu et al., Appl. Phys. Lett., 89, 062505 (2006).
[3] J.H. Lee et al., Appl. Phys. Lett., 90, 032504 (2007).
[4] H.K. Seong et al., Nano Lett., 7, 3366 (2007).
[5] P.R. Ganz et al., J. Cryst. Growth, in press: doi:10.1016/j.jcrysgro.2010.10.115
[6] Wilhelm et al., AIP Conf. Proc., 879, 1675 (2007).
[7] Y. Joly, Phys. Rev. B, 63, 125120 (2001).
__________________________
* Contact: ralf.schuber@kit.edu

MoP14
Growth Optimization for InAs/GaSb T2SL Structures by MBE
Y. X. Song 1* , S. M. Wang 1 , C. Asplund 2 , H. Malm 2 , X. Lu 3 , J. Shao 3
1
Department of Microtechnology and Nanoscience, Photonics Laboratory, Chalmers University of Technology,
SE-41296 Gothenburg, Sweden. 2 IRnova AB, Isafjordsgatan 22, SE-16440 Kista, Sweden. 3 National Laboratory
for Infrared Physics, Shanghai Institute of Technical Physics, CAS, Shanghai 200083, P. R. China
Molecular beam epitaxy (MBE) growth of InAs/GaSb type-II superlattices (T2SL) aiming for
mid-wavelength infrared (MWIR) photodetection was optimized in terms of interfacial quality, strain
compensation and photoluminescence (PL) intensity.
InAs/GaSb T2SL material is getting more intensively investigated in recently years due to its promising
applications in MWIR and long-wavelength infrared (LWIR) photodetections. Advantages, such as a
tailorable effective band gap, high quantum efficiency, potential high operation temperature, suitability
to fabricate focal plane arrays, etc. [1], make it very competitive compared with conventional infrared
detection materials such as mercury cadmium telluride (MCT), bulk InSb and quantum well infrared
photodetectors (QWIPs).
The interfacial quality was examined by X-ray diffraction (XRD) for T2SL samples grown at different
growth temperatures. The optimal growth temperature was found to be 340 °C judged by the narrowest
full width at half maximum (FWHM) of the superlattice XRD diffraction peaks. Sb and As soaking time
was then investigated to optimize the interfacial quality. The interface quality was found to be very
sensitive to the As soaking time and improved when changing from 0 s to 1s, but a longer As soaking
time has an adverse effect leading to possible strain relaxation due to the excess tensile strain brought by
the formation of GaAs-like interfaces. Strain compensation was tried first by increasing the Sb soaking
time to introduce InSb-like interfaces but it was not effective. Additional InSb layer was then introduced
to different positions of the T2SL structure for strain compensation. As shown in Fig.1, the 0 th order
diffraction peak in the XRD rocking curve moves toward the GaSb peak for the sample with 1ML InSb
insertion, indicating effective strain compensation. Insertion of an InSb layer was also found to improve
the interfacial quality. The optimized FWHM of the 1 st order diffraction peak reaches 39 arcsec, only
slightly larger than 25 arcsec of the GaSb substrate peak.
Fig. 2 shows PL spectra of three samples at 77K. The PL intensity is correlated to the interfacial quality.
The sample with 1 s As soaking (red) has smaller FWHM of the XRD diffraction peaks than those of the
sample with 0.5 s As soaking (blue), and the PL intensity is also stronger. The sample with 0.5 s As
soaking and an InSb compensation layer has the lowest FWHM value and shows the strongest PL
intensity. Moreover, insertion of an InSb layer red-shifts the PL peak wavelength.
Fig 1. XRD (004) rocking curves of samples with (blue)
and without (red) an InSb strain compensation layer.
[1] E. Plis and A. Khoshakhlagh, et al. J. Vac. Sci. Technol. B 28, 3 (2010)
_________________________
* Contact: yuxin.song@chalmers.se
Fig 2. PL spectra at 77K of samples with 0.5 s (blue &
green) and 1 s (red) As soaking, and with (green) and
without (blue and red) an InSb strain compensation layer.

MoP15
A prototype of heterovalent interfaces:
Reduction of the potential barrier in the conduction band at the
n-ZnSe / n-GaAs interface by Se predeposition
A. Frey, U. Bass, S. Mahapatra, C. Schumacher, J. Geurts and K. Brunner *
Physikalisches Institut (EP3), Universität Würzburg, Am Hubland, 97074 Würzburg, Germany
Heterovalent semiconductor heterointerfaces are unique in that they exhibit variable band offsets due
to the electric dipole moments associated with different atomic interface configurations [1, 2]. A
prototype of such heterovalent interfaces is that of ZnSe / GaAs, which also plays a role in dilute
magnetic II-VI / III-V semiconductor heterostructures for electron spin injection and in II-VI
optoelectronic devices.
In the presented work the electronic and structural properties of n-ZnSe / n-GaAs heterostructures
grown by molecular beam epitaxy with a fractional monolayer of Se predeposited at II-VI growth start
are studied by means of temperature-dependent electric transport across the heterointerface,
electrochemical capacitance-voltage (CV) profiling, Raman spectroscopy, high resolution x-ray
diffraction (HRXRD) and etch pit density (EPD) measurements [3].
By measuring the activation energy for thermally activated transport across the heterointerface
(Fig. 1), we find that the potential barrier in the conduction band at a Zn-rich n-ZnSe / n-GaAs
heterointerface is as high as 550 meV. With Se predeposition it gradually decreases down to about
70 meV, which entails a decrease of resistivity of the heterointerface by five orders of magnitude
(Fig. 2). The structural layer properties assessed by HRXRD and EPD for a conventional Zn-rich
growth start and for Se predeposition are comparable. It is remarkable that the layer thickness
decreases with increasing amount of predeposited Se by as much as 35 nm. This is assigned to a
significantly varying initial ZnSe growth rate depending on Zn or Se coverage of the As-terminated
GaAs(001) surface [4]. CV measurements show a large depletion region at the heterointerface with an
electron deficit of 1.5x10 13 cm -2 , which is independent of the growth start procedure. Raman
measurements of the relative intensity of longitudinal-optical (LO) phonon and coupled plasmon-LO
phonon modes of ZnSe and GaAs show, however, that Se predeposition partially shifts this depletion
region from GaAs into ZnSe, compared to Zn-predeposition [5].
The results are discussed in terms of a band-bending model accounting for a variable ZnSe / GaAs
band offset, acceptor-type interface states, and atomic intermixing profiles caused by Zn diffusion into
GaAs and As segregation into ZnSe, which all depend on the II-VI growth
start procedure (Fig. 3). From these we conclude that Se predeposition significantly lowers the CBO,
but the finite acceptor type interface states and the about 50 nm wide depletion region caused by
intermixing still limit the electronic conductivity. The reduction of the ZnSe/GaAs CBO with Se
predeposition observed in our electrical measurements agrees with first principles calculations of Kley
and Neugebauer [1], but appears to be in contrast to the results derived from photoelectron
spectroscopy of Nicolini et al. [2]. Possible influences of details of GaAs surface preparation and
growth conditions on the band alignment will be discussed.
[1] A. Kley and J. Neugebauer, Phys. Rev. B, 50, 8616 (2010).
[2] R. Nicolini, L. Vanzetti et al., Phys. Rev. Lett, 72, 294 (2010).
[3] A. Frey, U. Bass, S. Mahapatra, C. Schumacher, J. Geurts and K. Brunner, acc. for publ. in Phys. Rev. B, 82 (2010).
[4] A. Benkert, C. Schumacher, K. Brunner and R. Neder, Appl. Phys. Lett. 90, 162105 (2007).
[5] A. Frey, F. Lehmann, P. Grabs, C. Gould, G. Schmidt, K. Brunner and L. W. Molenkamp,
Semicond. Sci. Technol., 24, 35005 (2009)
__________________________
* Contact: brunner@physik.uni-wuerzburg.de

MoP15
Fig 1: Arrhenius plot of thermally activated electron transport across an n-ZnSe / n-GaAs heterointerface at three different
bias voltages. From the linear regime the interface potential barrier height Φ B is determined.
Fig 2: Current-voltage characteristics measured through several n-ZnSe / n-GaAs interfaces with gradually varied amounts of
Zn or Se predeposited before ZnSe growth start.
Fig 3: Modeled conduction band edge and electron density profiles for Zn- and Se-rich ZnSe/GaAs heterointerfaces.

MoP16
Effect of growth temperature on quantum dot laser (Ga,In) (N,As)
self-assembled quantum dots
O. A. Niasse 1 , M. AL Khalfioui 2 , B. Ba 1 , A. Bèye 3 , M. Leroux 2
1 Laboratoire des Semi-conducteurs et d'Energie Solaire, 3 Groupe Physique des Matériaux, Département de
Physique Faculté des Sciences et Techniques Université Cheikh Anta Diop, Dakar Sénégal
2 Centre de Recherche sur l’Hétéro-épitaxie et ses Applications du CNRS, 06560 VALBONNE FRANCE
ABSTRACT
In this paper, we have made quantum well laser structures and self assembled quantum dots on GaAs
substrate. The samples studies in this work were grown by molecular beam epitaxy (MBE) Laser
structures with InAs quantum dots encapsulated GINA are performed under different growth
conditions and compared with boxes wrapped with InGaAs. The influence of the encapsulation layer
GINA that moves the emission wavelength towards the red is highlighted. The incorporation of
indium and nitrogen was monitored and optimized by optical characterization. The growth temperature
is one of the growth parameters the most important and studied through images of nanostructures
obtained by SEM and AFM. Ohmic contacts deposited by electron evaporator or by Joule effect after
etching by photolithography have allowed us to complete the characterization by performing electrical
measurements that provide information on the sensitivity, stability, power consumption heating of the
structure built.
REFERENCES
[1] Wetting layer states of InAs/GaAs self-assembled quantum dot structures: Effect of intermixing and capping layer,
G. SJk,a_ K. Ryczko, M. Motyka, J. Andrzejewski, K. Wysocka, and J. Misiewicz, JOURNAL OF APPLIED PHYSICS
101, 063539 _2007
[2] Slow light control with electric fields in vertically coupled InGaAs/GaAs quantum dots
Chun-Hua Yuana_Ka-Di Zhub_ and Yi-Wen JiangJOURNAL OF APPLIED PHYSICS 102, 023109 2007
[3] Photoluminescence investigation of low-température capped self-assembled InAs/GaAs quantum dots,
R. Songmuang, S. Kiravittaya, M. Sawadsaringkan, S. Panyakeow, and O. G. Schmidt, J. Cryst. Growth 251, 166 (2003)
[6] Electroluminescence analysis of 1.3-1.5 μm InAs quantum dot LEDs with Ga,In) (N,As) capping layers, M. Montes1,*,
A. Hierro1, J. M. Ulloa1, A. Guzmán1, M. Al Khalfioui2, M. Hugues2,**, B. Damilano2, and J. Massies, P hys. Status Solidi
C 6, No. 6, 1424–1427 (2009)
[7] Photoluminescence of self-assembled InAs/GaAs quantum dots excited by ultraintensive femtosecond laser,
Shihua Huang1,a_ and Yan Ling2, JOURNAL OF APPLIED PHYSICS 106, 103522 2009

Intensité lumineuse en u.a
Intensité d'EL
MoP16
(a)
(b)
Figure1 : SEM and AFM images for different growth conditions
35 S1008
S1009
30
S 1005
S 997
25
20
15
10
10 -1 Courant (mA)
1090
10 -2
1000
800
500
200
10 -3
5
0
10 -4
-5
0,7 0,8 0,9 1,0 1,1 1,2 1,3 1,4 1,5 1,6
E(eV) 0,85 0,90 0,95 1,00 1,05 1,10
Energie d'EL (eV)
Fig. 2. PLE at E det equal to QD ground state and PL spectra at excitation density of 15
mw/cm2 measured at 10 K for
10 -5
a) sample A2, b) sample B3
Echantillons Tension de
seuil
Résistance
S997 0,54V 8.10 -2
QD
S1009 0,54V
S1008 0,5V
Tableau 1: Threshold voltage

MoP17
Effect of different monolayer coverage for the seed layer in
quaternary alloy capped multilayer InAs/GaAs quantum dot system
A. Mandal , N. Halder and S. Chakrabarti *
Centre for Nanoelectronics, Department of Electrical Engineering, Indian Institute of Technology Bombay,
Mumbai – 400076, Maharashtra, India.
Multilayer InAs/GaAs Quantum dots (QDs) are attractive material system for their possible use in the
active region of the third generation Intermediate band solar cells (IBSCs). InAs QDs in IBSCs can
extract a larger portion of the long wavelength region of solar radiation for photocurrent generation.
Multilayer (10 layer) of InAs dots were grown over semi insulating GaAs substrate (001) by solid
source molecular beam epitaxy (SSMBE). Firstly, an intrinsic GaAs buffer layer of thickness 0.4 µm
was grown at 590°C on the GaAs substrate. Then the substrate temperature was ramped down to
520°C during the growth of the 1000Å intrinsic GaAs layer. To grow self assembled seed layer of
InAs quantum dots (QDs), for one sample, the InAs monolayer (ML) coverage was 2.7 ML (Sample
A) while for the other one, it was 2.5 ML (Sample B). The seed layer in both the samples was capped
with a combination of 20Å quaternary capping of In 0.21 Al 0.21 Ga 0.58 As and 90Å intrinsic GaAs layer.
Later, 9 periods of 2.5 ML active InAs dots were grown at 480°C to avoid In/Ga intermixing for both
the samples with a capping combination of 20Å In 0.21 Al 0.21 Ga 0.58 As and 130Å intrinsic GaAs layer
(Figure 1a and 1b). The growth rate for GaAs layer, quaternary alloy and InAs dots were kept at 0.58
µm/h, 0.95 µm/h and 0.2011 ML/s respectively. It has already been proved that the key issue in bilayer
structure is the templeting effect of the seeded QD layer which dictates the upper layer’s QD number,
density and the stacking. In this article we have showed the effect similar to that of the bilayer one [1];
in MQD structures, the seed layer or the lower most layer controls the stacking of the QDs in the upper
layers. In sample A, greater ML coverage for the seed / lower most layer produce greater surface
strain; as a result the phase separation in quaternary alloy [2] was also increased which resulted in
more aggregation of the In atoms (from the alloy itself) near the elastically relaxed islands in Sample
A. This incident simultaneously reduced the out diffusion or segregation of In atoms from the QDs,
during the growth of the top layers. Further, higher monolayer coverage in the seed layer in sample A,
helps in good propagation of templeting effect in the upper layers of the MQD structure which helps in
vertical stacking of QDs. That is why, in the case of Sample A, we find more number of completed
stacks in comparison to Sample B, as seen from the cross sectional transmission electron microscopy
(XTEM) pictures (Figure 2a and 2b). 23% and 7.15% stoppage for the QD stacks, base width of 22 nm
and 28 nm, height of 9 nm and 10 nm were measured from XTEM results at the 10 th layer dot for
Sample A and B respectively. The lateral size of the dots in Sample B are slightly increased compared
to Sample A, this may be due to In/Al intermixing in the dots, due to lesser surface strain on the dots
[2]. Photoluminescence (PL) emission spectra showed that the overall integrated PL intensity of
Sample A is higher than that of Sample B due to the presence of greater amount of complete QD
stacks and more amount of radiative recombination centers. Moreover, from relative integrated PL
plots (Figure 3), activation energy E a was calculated to be 125.3 meV for Sample A while for Sample
B it was 67.6 meV. Increased activation energy proved better carrier confinement within Sample A
compared to Sample B. Thus, our study revealed that if greater strain is produced due to greater
monolayer coverage at the seed / lower most layer of MQD structures, it ultimately enhances the dot
quality in the QD stacks. DST, India is acknowledged.
[1] S. Chakrabarti, N. Halder, S. Sengupta, S. Ghosh, T. D. Mishima and C. R. Stanley, Nanotechnology, 19, 505704 (2008).
[2] S. Adhikary, N. Halder and S. Chakrabarti, Journal of Crystal Growth, 312, 724 (2010).
__________________________
* Contact: subho@ee.iitb.ac.in, subhanandachakrabarti@gmail.com

MoP17
Figures:
(a)
Fig. 1: Heterostructures for, (a) Sample A, (b) Sample B.
(b)
(a)
Fig. 2: XTEM images for, (a) Sample A, (b) Sample B.
(b)
Relative integrated PL intensity (A.U.)
1
0.36788
0.13534
0.04979
0.01832
0.00674
0.00248
Sample A, E a = 125.3 meV
Sample B, E a = 67.6 meV
0.00 0.02 0.04 0.06 0.08 0.10 0.12 0.14
1/ T (1/Kelvin)
Fig. 3: Activation energy calculation from relative integrated PL plots for Sample A and Sample B.

MoP18
Neutron reflectometry studies of hetero-interfacial H layer in
highly lattice-mismatched epitaxy on Si
H. Asaoka*, T. Yamazaki, D. Yamazaki, M. Takeda and S. Shamoto
Japan Atomic Energy Agency (JAEA), Tokai, Ibaraki 319-1195,Japan
Hetero-structures with atomic order thickness are fabricated successfully by using
molecular beam epitaxy (MBE) methods. Hetero-epitaxial growth, however, is possible
only for limited material combinations resulting from their lattice mismatches
Because of the H mono-atomic layer on the Si surface, Sr or SrO layer has grown
epitaxially in spite of the large lattice mismatch [1,2]. It could be interesting to see
how the interfacial mono-atomic layer can accommodate the lattice mismatch.
Neutron reflectometry is one of the best techniques to reveal the role of the H layer in
crystal growth.
Fig. 1(a) shows the reflectivity profile of chemically treated mono-atomic H- or D-Si
(111) substrate before Sr deposition using SUIREN at JAEA. The difference of both
surfaces appears in the reflectivity profiles in the vicinity of 0.09 Å -1 . The detail
reflectivity measurements (Fig. 1(b)) reveal the mono-atomic H or D layer on Si
surface is clearly distinguishable by using the neutron reflectometry.
Strontium layers grow on the mono-atomic H- or D-Si surface. The difference of both
samples appears in the reflectivity profiles resulting from a contrast variation between
H and D atoms at the buried hetero-interface. The results suggest the existence of H layer at
the hetero-interface acting as an effective buffer layer to manage the highly mismatched epitaxy on
Si.
Fig. 1: (a) Neutron reflectivity profile of chemically treated mono-atomic H- or D-Si (111) substrate, and (b) detail
reflectivity profile in the vicinity of 0.09 Å -1 with accumulated measurement period which is fifth times as long as the
period of wide Q data of (a).
[1] H. Asaoka, K Saiki, A Koma and H Yamamoto, Thin Solid Films, 369 273 (2000)
[2] H. Asaoka, T. Yamazaki and S. Shamoto, Appl. Phys. Lett., 88 201911 (2006)
__________________________
* Contact: asaoka.hidehito@jaea.go.jp

MoP19
Surface Electronic Properties of GaAs Nanowires.
O. Demichel 1,2 , M. Heiss 1 , J. Bleuse 2 , H. Mariette 2 , A. Fontcuberta 1
1
EPFL, Lausanne, Switzerland.
2
CEA-Grenoble, 17, rue des Martyrs, F-38054, Grenoble, France.
ABSTRACT BODY:
Semiconducting nanowires (NWs) are a topic of intense research as they offer the opportunity to
explore properties of one-dimensional electronic systems. Their electronic properties are promising for
applications in nanoelectronics or photodetection as well as for sensing applications. The understanding and the
mastering of the electronic properties of such one-dimensional systems are essential to achieve high efficiency
devices. In particular, with the increase of the surface/volume ratio, the electronic properties of NW based
devices become strongly dependent on the surface electronic states. Indeed, surfaces electronic states can appear
in the electronic band gap, altering the NW electronic properties. Two main effects are reported in literature:
surface states can act as recombination centers for free carriers or as surface charged traps. We has recently
quantified this first effect by an original optical method for silicon NWs [1-2]. The surface charge traps induce a
pinning of the Fermi-level at the surface and a depletion shell appears: the electronic canal available for carriers
is decreased.
In this letter, we present our results on the influence of {110} surfaces on GaAs NWs measured by low
temperature micro-photoluminescence (μ-PL) spectroscopy. We compare unpassivated NWs with those which
were capped (passivated) with a shell of Al 0.4 Ga 0.6 As. NWs were obtained by molecular beam epitaxy have a
prismatic cross section [3]. Moreover, we take advantage of the tapered shape of the NWs to understand the role
of diameter and surface/volume ratios in the luminescence efficiency. We demonstrate that capping directly
modifies the recombination velocity and passivated NW electronic properties are governed by surface
recombinations whereas unpassivated NWs are strongly depleted by the Fermi-level pinning at the surface
induced by charged surface trap states. Finally, we measured a surface recombination velocity of 3.10 3 cm.s -1 for
passivated NWs: one order of magnitude lower than values previously reported for {110} GaAs surfaces. And
the surface charged trap density of uncapped NWs (~10 12 cm -2 ) is in agreement with values of surface charged
trap density of native oxidized GaAs surfaces. Photo-courant measurements were then performed and we
demonstrate the increase of the photo-voltage conversion efficiency once NWs are passivated. This work will
serve as a guidance for the passivation of the NW surfaces which is of crucial importance for applications in
laser and solar cell technology.
[1] O. Demichel, V. Calvo, N. Pauc, A. Besson, P. Noé, F. Oehler, P. Gentile, and N. Magnea, Nano Lett. 9 (7),
2575 (2009).
[2] O. Demichel, V. Calvo, A. Besson, P. Noé, B. Salem, N. Pauc, F. Oehler, P. Gentile, and N. Magnea, Nano
Lett. 10 (7), 2323 (2010).
[3] A. Fontcuberta i Morral, D. Spirkoska, J. Arbiol, M. Heigoldt, J. R. Morante, and G. Abstreiter, Small 4 (7),
899 (2008), ISSN 1613-6829.

MoP20
Critical thickness of 2D-3D and “hut”-“dome” transitions
at the growth of Ge x Si 1-x and Ge/Ge x Si 1-x layers on the Si(100)
V.A. Timofeev * , A.I. Nikiforov , V.V. Ulyanov , O.P. Pchelyakov
Rzhanov Institute of Semiconductor Physics SB RAS,
Lavrentjeva 13, 630090 Novosibirsk, Russia
The silicon structures with germanium quantum dots are of practical interest for optoelectronics due to
their potential covering the regions from IR through the wavelengths used in fiber-optic
communications. Reflection high-energy electron diffraction (RHEED) is the most used technique in
MBE. There are available numerous papers that report studies of early stages of Ge growth on the
Si(100) surface but only few data on the influence of Ge x Si 1-x layer on the wetting layer thickness and
“hut”-“dome” transition [1].
A Katun-C MBE installation equipped with two electron beam evaporators for Si and Ge was used for
synthesis. As the deposited layer increases in thickness, elastic strains induced by mismatching of the
Si and Ge lattice constants also increase. Starting with some critical thickness, from two-dimensional
to three- dimensional growth mechanism (2D-3D) and from “hut”-clusters to “dome”-clusters (hutdome)
transitions is observed, a part of strains being relaxed that is energetically favorable due to a
decrease in the free energy of the system. Thus, identifying the moment of 2D-3D and “hut”-“dome”
transitions at various thickness of Ge x Si 1-x layer at 500°C allowed the 2D-3D and “hut”-“dome”
transitions thickness of Ge film to be determined as a function of Ge x Si 1-x thickness for different Ge
content in Ge x Si 1-x layers (see Figure 1a) and the 2D-3D transitions thickness of Ge x Si 1-x layer as a
function of Ge content in Ge x Si 1-x layers (see Figure 1b).
critical thickness of Ge (nm)
1.2
1.0
0.8
0.6
0.4
Ge / Ge x
Si 1-x
/ Si(100)
T s =500 o
x=0.15
x=0.3
x=0.6
Huts - Domes
x=0.15
x=0.3
x=0.6
0.2
2D-3D
0.0
0 2 4 6 8 10 12
thickness of Ge x
Si 1-x
(nm)
(a)
10 3
10 2
10 1
550 o C
2D-3D
10 0
750 o C
500 o C
0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0
Ge content (x)
Fig. 1: Critical thicknesses, (a) 2D - 3D and “hut” - “dome” transitions thickness of Ge film as a function of Ge x Si 1-x
thickness for different Ge content in Ge x Si 1-x layers, (b) 2D - 3D transitions thickness of Ge x Si 1-x layer as a function of Ge
content in Ge x Si 1-x layers.
critical thickness of Ge x
Si 1-x
layer (nm)
(b)
d D-Т
d М-B
d exp
The critical transition thicknesses are observed to decrease to reach saturation as the solid solution
layer thickens or its content increases. The observed decrease in thicknesses is accounted for by
strengthening the strain deformation in the solid solution layer. Theoretical dependences of critical
thicknesses from composition reflecting the appearance misfit dislocations at the substrate
temperatures of 550 о С (Dodson и Tsao) and 750 о С (Matthews и Blakesley) are shown at the figure 1b
for comparison [2]. The received dependence of the 2D-3D transition ( figure 1b) will enable to obtain
two-dimensional dislocation-free Ge x Si 1-x layers, which are used at the growth of multilayer periodical
Ge/Si sructures.
[1] Zhensheng Tao et.al. Applied Surface Science, 255 3548 (2009).
[2] K. Brunner. Rep. Prog. Phys., 65, 27-72 (2002).
__________________________
* Contact: Vyacheslav.t@isp.nsc.ru

MoP21
Epitaxy and characterization of GaMnAs
Martin Utz 1 , D. Schuh 1 , D. Bougeard 1 and W. Wegscheider 2
1
Universität Regensburg, Institut für Angewandte und Experimentelle Physik,
93053 Regensburg, Germany
2 ETH Zürich , Laboratorium für Festkörperphysik,
8093 Zürich, Switzerland
The dilute magnetic semiconductor GaMnAs is a prominent candidate for spin injection experiments into
III-V semiconductor hetero-structures. Here, we report on our work to realize GaMnAs films with high
structural quality. The growth was especially focused on two issues:
On the one hand side we focused on films with high Mn concentration in order to maximize the Curie
temperature. The second focus lay on thin layers which are currently used to study the coupling of iron and
GaMnAs in hybrid structures [1].
All presented layers were grown on semi-insulating (100) GaAs substrates and a 300nm high temperature
GaAs and 5 to 10nm low temperature GaAs buffer.
The surface quality was monitored in situ by RHEED and was checked post-growth with AFM in air. Both
– annealed and as grown samples - have been characterized:
The Curie temperature was determined by measuring the point singularity in the temperature derivative of
the resistivity [2]. Additionally, charge carrier density and mobility were determined by Hall measurements
at room temperature in the paramagnetic state of the samples.
For the enhancement of T C GaMnAs layers of thicknesses from 18nm to 50nm and Mn contents up to
13,7% have been grown, exceeding the well established 6% used in many structures (Fig. 1). On that
purpose the substrate temperatures - measured by band edge spectroscopy – were reduced down to
temperatures as low as 120°C as well as the beam equivalent pressure ratios (As 4 : Ga) from 3 to 1.
In this growth series Curie temperatures up to 173°C could be achieved (Fig. 2). The mobility of the charge
carriers is increasing along with the impurity concentration, despite the absence of ferromagnetism at room
temperature. In contrast, the charge carrier density is slightly decreasing.
Thin GaMnAs layers have been grown in the range of 3nm to 18nm – all samples with a Mn content of 6%.
Down to 10nm the samples still showed bulk properties in Curie temperature, charge carrier density and
mobility. Below, distinct changes in magnetic properties take place. Interestingly, below 10 nm thickness
we also observe a changing in the coupling mechanism to adjacent Fe layers.
Fig 1: Overview of the Curie temperatures of the grown
samples with thickness ≥ 18nm
Fig 2: temperature dependence of the resistivity of an
annealed sample (13,7% Mn) and its derivative
[1] Matthias Sperl et al., PRB, 81, 035211 (2010)
[2] Vit Novak et al., Phys. Rev. Lett., 101, 077201 (2008)

MoP22
Mid-infrared Quantum Dot LEDs and microdisk laser grown by MBE
A. Hochreiner 1 , M. Eibelhuber 1 , T. Schwarzl 1 , H. Groiss 1 , V. Kolkovsky 2 ,
G. Karczewski 2 , T. Wojtowicz 2 , W. Heiss 1 , G. Springholz 1
1 Institute of Semiconductor Physics, University of Linz, Austria
2 Polish Academy of Sciences, Warszawa, Poland
Self-assembled semiconductor quantum dots have attracted tremendous interest as active medium in
optoelectronic devices. For mid-infrared (MIR) devices narrow gap IV-VI semiconductor compounds
are well suited because of their favorable electronic band structure and low Auger recombination rates.
However, conventional IV-VI Stranski-Krastanow QDs exhibit only weak luminescence due to the
strain induced change of the band alignements between the dots and the surrounding barrier material,
resulting in an unfavorable type II alignement [1]. For the realization of quantum dots with strong
infrared emission, we have therefore developed an alternative strain-free synthesis method in which
dot formation is induced by phase separation rather than by heteroepitaxial strain. The resulting QDs
exhibit almost spherical shapes with abrupt interfaces, are essentially defect- and strain-free and show
intense mid-infrared photoluminescence (PL) even at room temperature [2].
In this work, we show the first MIR device applications based on these unique PbTe QDs, epitaxially
embedded in wide band gap CdTe. In particular, we demonstrate the fabrication of mid-infrared light
emitting diodes (LEDs) operating in cw up to room temperature and we demonstrate cw optically
pumped lasing up to 200 K in microdisk structures.
The samples were grown by molecular beam epitaxy (MBE) on high quality CdTe buffer layers
predeposited on GaAs (001) substrates, with CdTe growth temperatures of 360°C and beam fluxes of
1 ML/s. The QD diode structures consist of a 0.5 µm intrinsic CdTe zone with the active PbTe dot
layer in the middle grown on a 4 µm thick n-doped CdTe buffer layer, followed by a 0.5 µm thick p-
doped CdZnTe cap layer with Zn content of 15%, as shown schematically in Fig. 1(b). For n- and p-
doping, iodine and nitrogen were used, respectively, with carrier concentration of about 10 18 cm -3 in
both cases. Typical growth temperatures for the PbTe layers are 265°C with growth rates of around 0.4
ML/s. Subsequent annealing leads to the formation of isolated PbTe QDs by phase separation as
illustrated by plan-view transmission electron microscopy (TEM) image shown in Fig. 1(a). Cw
electroluminesensce (EL) spectra were measured at various temperatures up to 300 K for dots with
average diameters of 10 and 12 nm, respectively, as shown in Fig. 1(c)-(d). The LED emission was
compared to PL spectra from the same sample region using a 1064 nm laser with photon energy well
below the CdTe band gap. At all temperatures from 30 K to 300 K (see Fig. 1), the EL exactly
matches the PL for both samples and even the spectral shape agrees remarkably well, proving that the
EL indeed arises from the embedded PbTe QDs. The electroluminescence of the 10 nm dot LED as a
function of injected current varying from 8 mA to 17 mA at 30 K is shown in Fig. 2. The integrated
output power increases linearly with diode current as is typical for spontaneous emission. A similar
behavior was found at higher operation temperatures. At 300 K the total output power was found to be
0.7 µW at 8 mA diode current [3].
To obtain lasing from the PbTe QDs, microdisks with a diameter of 40 µm were fabricated by
photolithography and wet chemical etching. The single active PbTe dot layer (see Fig. 3(c)) is
positioned in the center of the 2 µm thick CdTe waveguide. A SEM image of a single microdisk is
shown in Fig. 3(d). The QD microdisks were optically excited in cw below the CdTe band gap at a
wavelength of 1030 nm, resulting in laser emission up to temperatures as high as 200 K (Fig. 3(a))
with a maximum output power at 50 K of 0.15 mW considering homogenous emission. The laser
intensity at 200 K as a function of the pump power is depicted in Fig. 3(b). Below threshold, the
intensity is almost zero and above it increases linearly as expected for lasing. Thus, our unique PbTe
QDs proof their suitability for novel mid-infrared optoelectronic devices.
[1] M. Simma, et al., Appl. Phys. Lett. 88, 201105 (2006).
[2] W. Heiss, et al., Appl. Phys. Lett. 88, 192109 (2006); H. Groiss, et al., Appl. Phys. Lett. 91, 222106 (2007).
[3] A. Hochreiner, et al., submitted to Appl. Phys. Lett.
* Contact: astrid.hochreiner@jku.at

MoP22
(b)
(a)
100 nm
(b)
p-CdZnTe:N
PbTe/CdDs
i-CdTe
n-CdTe:I
GaAs (100)
Normalized intensity (arb. units)
EL
PL
120 K
200 K
300 K
(c) (a)
10 nm dots
1.8 2.4 3.0 3.6
Wavelength (µm)
12 nm dots
1.6 2.4 3.2 4.0
Wavelength (µm)
Fig. 1: (a) Plan-view dark-field TEM image of 10 nm PbTe quantum dots in CdTe formed after the post-growth
annealing. (b) Schematic illustration of the LED structure in which the dots are embedded in the middle of the
intrinsic zone of a CdTe/CdZnTe p-i-n junction. (c) Temperature dependent electro- and photoluminescence spectra
of PbTe/CdTe quantum dot diode mesas with different average dot size of 10 and 12 nm in (c) and (d), respectively.
The photoluminescence was excited with pump photon energies (1064 nm wavelength) below the CdTe band gap. For
both samples, a good agreement between the EL and PL spectra is found at all temperatures.
30 K
EL
PL
(d) (b)
140 K
190 K
230 K
300 K
EL Intensity (rel. units)
17 mA
15 mA
13 mA
10 mA
8 mA
T=30 K
17 mA
8 mA
output vs.
current
5 10 15 20
Current (mA)
2.0 2.5 3.0 3.5
Wavelength (µm)
( )
Fig. 2: Electroluminescence (EL) spectra of the PbTe/CdTe light emitting diode with 10 nm PbTe dots in the active
region measured at 30 K for different injection currents increasing from 8 to 17 mA. Inset: Integrated output intensity
versus injection current, showing a linear behaviour as indicated by the solid line.
Intensity (rel. u.)
10nmQDs
(d)
Fig. 3: (a) Microdisk laser spectra and (b) Intensity versus effective pump power at 200 K. (c) X-TEM image of PbTe
dots as in the active region. (d) SEM image of a microdisk with an active PbTe dot layer and a disk diameter of 40
µm. The undercut of GaAs is achieved by selective isotropic wet chemical etching and provides an effective mode
confinement in the vertical direction.

MoP23
High-quality structures of InAs QDs in Al 0.9 Ga 0.1 As matrix grown
by droplet epitaxy
A. A. Lyamkina 1,2 *, D. S. Abramkin 1 , D. V. Dmitriev 1 , S. P. Moshchenko 1 ,
T. S. Shamirzaev 1 , A. I. Toropov 1 , K. S. Zhuravlev 1
1
Rzhanov Institute of Semiconductor Physics, Lavrent’eva ave. 13, 630090, Novosibirsk, Russia
2
Novosibirsk State University, Pirogova 2, 630090, Novosibirsk, Russia
Semiconductor InAs quantum dots (QDs) are intensively investigated due to their perspective
application for fabrication of novel devices such as light emitting diodes. An important advantage of
InAs/AlAs QDs system is shifting of an emission peak to visible spectral range. However, as it was
shown in our recent works there is problem to produce effective light-emitting devices based on
InAs/AlAs QDs due to insufficient capture of charge carriers from wetting layer to QDs [1]. This
problem is caused by a heterointerface roughness that complicates a motion of charge carriers in the
wetting layer and high concentration of non-radiative recombination centers localized in the wetting
layer that capture charge carriers impeding their penetration into QDs. In this work we demonstrate
that perfect InAs/Al 0.9 Ga 0.1 As QDs structures with significantly flatter heterointerface and low
concentration of non-radiative recombination centers in the wetting layer can be grown by droplet
epitaxy.
The structures with InAs/Al(Ga)As QDs studied in this work were grown by molecular beam epitaxy
using a Riber32P system with gate arsenic source on semi-insulating (001)-oriented GaAs substrate.
The quality of the surface was controlled by means of reflection high energy electron diffraction
(RHEED). The diffraction patterns clearly demonstrate (5x2) surface reconstruction that corresponds
to metal-enriched surface structure [2]. Sharp diffraction spots in RHEED pattern indicate the presence
of large coherent terraces. The estimations based on an obtained longitudinal width of spots reveal an
average terrace size to be about 100 nm. In addition a surface prepared for QD growth was
investigated by AFM in air. AFM images also demonstrate atomic steps with an average terrace size
about 100 nm.
To fabricate QDs droplet epitaxy method was used. Two monolayers of indium were deposited to the
substrate at 510º С in the absence of arsenic flux and the array of initial In droplets formed. Then the
sample was exposed to a beam equivalent pressure (BEP) of 1.5∙10 5 Torr of As 4 and InAs quantum
dots crystallized. The growth was interrupted for 60 sec after In valve had been closed, and then QDs
were capped with 25 nm of Al 0.9 Ga 0.1 As.
PL spectra at the temperature of 4.2 K were measured at the range of 550 to 850 nm. The
luminescence of both the QDs and wetting layer (WL) was observed. The PL of WL contained several
well-defined peaks of phonon replicas that demonstrate high quality of its heterointerfaces. Since
phonon replicas in InAs/AlAs WL spectrum were previously observed as broad peaks only [3],
separate lines of phonon replicas in spectra of our samples mean significant improvement of interface
quality. Domination of the QD band in the PL spectrum and temperature dependence of QD and WL
integrated intensities confirm that carriers captured from matrix in the wetting layer than are
efficiently collected to QDs.
This work was supported by the RFBR (Project 10-02-00513) and and the Program of Fundamental
Studies of the Presidium of the RAS No. 27.
[1]T.S. Shamirzaev et al, Nanotechnology 21, 155703 (2010).
[2] A.M. Dabiran, P.I. Cohen. Journal of Crystal Growth 150, 23-27 (1995).
[3] T.S. Shamirzaev et al. Physical Review B 76, 155309 (2007).
_______________________________
*
Contact: lyamkina@thermo.isp.nsc.ru

MoP24
Structural properties of InAlN single layers nearly latticematched
to GaN grown by plasma assisted molecular beam epitaxy
Ž. Gačević 1,* , S. Fernández-Garrido 1,** and E. Calleja 1
D. Hosseini 2 , S. Estradé 2, 3 and F. Peiró 2
1
ISOM, Universidad Politécnica de Madrid, Avda. Complutense s/n, 28040 Madrid, Spain
2
LENS-MIND-IN2UB, Departament d’Electrònica, Universitat de Barcelona, Martí i Franquès 1, 08028
Barcelona, Spain
3
TEM-MAT, SCT- UB, Solé i Sabarís 1, 08028 Barcelona, Spain
The high lattice mismatch between III-nitride binaries (InN, GaN and AlN) remains a key
problem to grow high quality III-nitride heterostructures. Recent interest has been focused on the
growth of high-quality InAlN layers, with approximately 18% of indium incorporation, in-plane
lattice-matched (LM) to GaN. While a lot of work has been done by metal-organic vapour phase
epitaxy (MOVPE) by Carlin and co-workers, its growth by molecular beam epitaxy (MBE) is still in
infancy.
In our recent publications we reported on: the growth map for InAlN single layers and the
growth of 10-period LM InAlN/GaN DBRs [1-3] with highest reported reflectivity (~60%, Fig. 1),
within MBE growth technique. However, successful growth of In 0.18 Al 0.82 N/GaN DBRs is followed by
two difficulties, first, to obtain flat InAlN layers with good homogeneity and high crystalline quality
and, second, to find a compromise between different growth temperature ranges required for a good
In 0.18 Al 0.82 N (~535 ºC) and GaN (~700 ºC) growth. This work addresses the first commented issue, i. e.
the properties of InAlN nearly LM to GaN, grown by MBE.
The two samples under study (S1 and S2) are grown on c-plane GaN-on-sapphire templates,
under exactly same growth conditions (at 535 ºC, under effective stoichiometry, [1-3]) apart from their
growth times, being their thicknesses approximately 70 and 420 nm, respectively.
Figure 2 shows ω/2Θ and ω-rocking scans around [0002] Bragg spot. Note that there is no
indication of crystalline InN or AlN. The reciprocal space maps (RSMs) around symmetric [0002] and
asymmetric [10-15] reflections (Fig. 3) reveal that the samples have grown pseudo-morphically on the
underneath GaN, i. e. with the same in-plane lattice constant. Making use of linear approximation of
the III-nitrides elasticity theory we estimate the indium content to be around 19%. Note that the full
width of half maximum (FWHM) of ω-rocking curves of the InAlN and the underneath GaN are
practically equal: 0.1º/0.11º, for [0002], and 0.58º/0.55º, for [10-15] reflections, respectively. These
results confirm excellent crystalline quality of the InAlN layers. In addition, there is no evidence of
structural deterioration of the material with the layers’ increasing thickness.
Figure 4 features tapping-mode AFM measurements of the samples comparing them to a
reference GaN template. Root mean square (RMS) surface roughnesses, over 2.52.5 µm 2 scan area,
are found to be: 0.41, 0.59 and 0.61 nm for the GaN, S1 and S2 sample, respectively. While we
appreciate certain increase in RMS (from GaN to InAlN), there is no evidence of progressive surface
roughening with increasing InAlN thickness.
The InAlN/GaN DBRs exhibit much inferior refractive index contrast when compared to their
AlN/GaN counterparts (8% and 16%, respectively). This difference provokes much deeper penetration
of the incident light into the former structures and consequently, their higher sensitivity to the presence
of residual absorption (i. e. absorption at lower-than-band-gap energies). To estimate optical properties
of the layers, we performed absorption measurements. The residual absorption was found to be as high
as 1.6% and 14%, at 600 nm, in S1 and S2 samples, respectively. To determine actual origin of this
unwanted effect, our samples are currently under thorough transmitting electron microscopy studies.
[1] S. Fernández-Garrido, Ž. Gačević and E. Calleja, Appl. Phys. Lett 93, 161907 (2008)
[2] Ž. Gačević, S. Fernández-Garrido, and E. Calleja, Phys. Stat. Sol (c) Vol. 6, S643-S645 (2009)
[3] Ž. Gačević, S. Fernández-Garrido, D. Hosseini, S. Estradé, F. Peiró and E. Calleja, accepted for publication in JAP (2010)
__________________________
*
Contact: gacevic@die.upm.es
**
Present address: Paul-Drude-Institut for Solid State Electronics, Hausvogteiplatz 5-7, 10117 Berlin, Germany

MoP24
Diffraction intensity (a. u.)
XRD around
[0002] Bragg spot
GaN
(FWHM ~0.11º)
InN
S2. GaN ω-rock
S2. InAlN ω-rock
S2. ω/2Θ
S1. GaN ω-rock
S1. InAlN ω-rock
S1. ω/2Θ
In (FWHM ~0.1º)
0.19
Al 0.81
N
AlN
Fig 1: State of the art of InAlN/GaN DBRs grown by
either MBE or MOVPE.
15 16 17 18 19
ω (º)
Fig 2: ω/2Θ and ω-rocking scans around [0002] Bragg
spot. The scans are characterized by no trace of phase
separation and good FWHM of InAlN ω-rocking curves.
Qy*10000(rlu)
Qy*10000(rlu)
3035
7605
InAlN
7505
InAlN
2995
7405
GaN
2955
GaN
7305
2915
-20 0 20 40
Qx*10000(rlu)
7205
2595 2795 2995
Qx*10000(rlu)
Fig 3: RSMs around symmetric [0002] (left) and asymmetric [10-15] (right) reflections (sample S2). (Left)
ω/2Θ scan (Fig. 2) corresponds to the scan along Q x =0 line, whereas GaN and InAlN ω-rocking scans (Fig.
2) correspond to Q y *10000 = 2971 and 3007 lines. (Right) The InAlN layer is in-plane LM to GaN.
S1
S2
80
S2 - Thick InAlN
S1 - Thin InAlN
Template GaN
GaN
Absorption (%)
40
0
300 400 500 600
Wavelength (nm)
Fig 4: 2.52.5 µm 2 AFM images of the samples under study.
Fig 5: Absorption measurements of the samples under study.
The onset of GaN absorption is marked by the arrow.

MoP25
Composition studies of site-controlled quantum dots
G. Biasiol 1,* , V. Baranwal 1 , S. Heun 2 , M. Prasciolu 1 , M. Tormen 1 , A. Locatelli 3 , T.
O. Mentes 3 , M. N. Orti 3 , and L. Sorba 2
1
Istituto Officina dei Materiali CNR, Laboratorio TASC, I-34149 Trieste, Italy
2 NEST Istituto Nanoscienze-CNR and Scuola Normale Superiore, I-56127 Pisa, Italy
3 Sincrotrone Trieste S.C.p.A., I-34149 Trieste, Italy
Despite the remarkable new physics phenomena associated to self-assembled quantum dots (QD), the
lack of spatial and spectral control due to the intrinsic randomness of nucleation in the Stranski-
Krastanow growth mode strongly limits their device application. To overcome this drawback, many
efforts have been devoted to reach a deterministic nucleation control of QDs. The most promising
approach to realize such controlled nucleation involves a pre-patterning the sample surface with an
array of holes, which act as preferential nucleation sites for the QDs [1]. To improve the spectral
uniformity of QDs through this site-control technique, a higher homogeneity must be reached both in
terms of size and composition, with respect to standard SK dots. A reduction of height fluctuation from
about 12% to about 7% has been obtained for QD stacks grown on hole arrays defined by electronbeam
lithography (EBL), with a substantial narrowing of the photoluminescence linewidth [2].
In this work, we analyze both the size and the composition distribution of site-controlled InAs/GaAs
QDs, and compare them with those obtained for QDs formed on planar surfaces. Two dimensional
morphological and composition maps were obtained on the same samples by Atomic Force
Microscopy (AFM) and X-ray Photoemission Electron Microscopy (XPEEM), respectively. Optical
quality of the samples was checked by low-temperature (4K) Photoluminescence (PL). Samples
containing 5 QD stacks and separated by 10nm GaAs were grown by MBE on (001) GaAs substrates
pre-patterned with square arrays of holes with 80nm diameter and periods from 250 to 600nm (growth
and fabrication details can be found in Ref. [3]). Surface concentration maps were obtained at the
Nanospectroscopy beamline at Elettra, Trieste [4].
Fig. 1 (left) shows an AFM image of an area of the sample patterned with a 250 nm period. Although
most of the holes are filled with a single QD, hole filling is not complete, and some holes with a
double dot are present as well. Consistently with what observed in Ref. [5], dots nucleate on the sides
of the holes, which are elongated in the [110] direction. No dots were observed between the holes. Fig.
1 (bottom) shows an In concentration map from the same 250 nm period area. QDs are visible as
higher-In content spots, with respect to the surrounding wetting layer. Note that the measured In
concentration (around 0.9) far exceeds the one generally reported in cross-sectional analysis of QD
composition (see, e.g., Ref. [6]). This is due to In surface segregation, together with the high surface
sensitivity of the XPEEM technique (about 0.5nm) [4]. Similar topographic and composition maps
were obtained for larger periods of the pattern, and for dots nucleated on unpatterned areas. We have
performed a statistical analysis of height and composition for dots grown on the different patterns, and
on the planar surface. We did not observe any trend with the period of the patterns; besides no
difference was detected between dots completely surrounded by filled holes and dots adjacent to voids.
Fig. 2a and b show histograms of the height distribution for single dots grown on all the periods and
on the planar surface, respectively. Average heights resulted to be 7.6±0.7 nm and 7.1±1.0 nm,
respectively. Fig. 2c and d show histograms of the In composition distribution for the same dots. The
average values were 0.917±0.004 and 0.933±0.008 inside and outside the patterns, respectively. This
analysis shows that growth of QDs on patterned templates results in a slight size increase and higher
material intermixing, and in a uniformity improvement of about 30% in height and 50% in
composition, with respect to standard SK growth. PL experiments on an identical sample capped with
100nm GaAs confirmed the similarity of dots grown on the different periods, with energy shifts of the
order of 10meV in going from 250 nm to 600 nm.

MoP25
Fig 1: Left: AFM image of an InAs/GaAs QDs array grown by MBE on a 250-nm period square hole array defined by
electron beam lithography. Right: 2D In composition map of the same pattern (different area) obtained by XPEEM.
Fig 2: a) Height histogram of site-controlled QDs grown on hole arrays with periods ranging from 250 to 600 nm. c)
In composition histogram for dots from the same arrays. b) height and d) composition histograms for dots grown on
planar surfaces.
__________________________
* Contact: biasiol@tasc.infm.it
[1] H. Heidemeyer, U. Denker, C. Müller, and O. G. Schmidt, Phys. Rev. Lett. 91, 196103 (2003).
[2] S. Kiravittaya, A. Rastelli, and O. G. Schmidt, Appl. Phys. Lett. 88, 043112 (2006).
[3] G. Biasiol, V. Baranwal, S. Heun, M. Prasciolu, M. Tormen, A. Locatelli, T. O. Mentes, M. A.
Niño, and L. Sorba, J. Cryst. Growth (in press).
[4] G. Biasiol, S. Heun, G. B. Golinelli, A. Locatelli, T. O. Mentes, F. Z. Guo, C. Hofer, C. Teichert,
and L. Sorba, Appl. Phys. Lett. 87, 223106 (2005).
[5] P. Atkinson, S. Kiravittaya, M. Benyoucef, A. Rastelli, and O. G. Schmidt, Appl. Phys. Lett. 93,
101908 (2008).
[6] A. Rosenauer, D. Gerthsen, D. Van Dyck, M. Arzberger, G. Boehm, and G. Abstreiter, Phys. Rev.
B 64, 245334 (2001).

MoP26
Single InAs quantum dots morphology and local electronic
properties on (113)B InP substrate
C. Cornet 1,* , P. Turban 2 , N. Bertru 1 , S. Tricot 2 , O. Dehaese 1 and A. Le Corre 1
1
Université Européenne de Bretagne, France
INSA, FOTON, UMR 6082, F-35708 RENNES
2 Equipe de Physique des Surfaces et Interfaces, Institut de Physique de Rennes UMR UR1-CNRS 6251,
Université de Rennes 1, F-35042 Rennes Cedex, France
Recent research developments have focused on the ability to control the crystalline growth at
the nanoscale. These research efforts have led, among other successes, to the fabrication of InAs/GaAs
Quantum Dots (QDs) which operate at the telecommunication wavelength of 1.3 µm.[1] Large efforts
to push the InGaAs/GaAs system to 1.55 µm wavelength have been hampered by the large strain that
accumulates in- and outside the QD structure during the Stranski-Krastanow growth leading to the
formation of plastically relaxed QDs. InAs QD growth on InP substrates has been proposed to
overcome this problem. However the growth of InAs on InP (100) leads to the undesirable coexistence
of nanostructures presenting various shapes, like quantum wires, quantum dashes or QDs.[2] Growth
on (113)B substrate has shown large improvements of QDs structural properties as well as promising
device performances.[3] Such substrate orientation is expected to have strong impact on QD shape as
demonstrated on InAs/GaAs nanostructures.[4] Here, we propose a characterization at the atomic scale
of InAs/InP (113)B QDs morphology by Scanning Tunneling Microscopy (STM) as well as the
determination of local electronic properties by Ballistic Electron Emission Microscopy (BEEM).
Samples presented here have been grown on doped InP:n (113)B substrate using gas source
MBE. After a 400 nm InP:n buffer layer, a 90 nm lattice-matched n-doped In 0.8 Ga 0.2 As 0.435 P 0.565 layer
has been deposited. Such quaternary alloy is used as optical confinement layer in laser guides [3].
After a 10 nm undoped InGaAsP layer, and a growth interruption (GI) under AsH 3 and PH 3 , 2.1 ML
InAs was deposited at 480 °C, with subsequent 30 seconds As GI. Samples were cooling down at the
room temperature, and an amorphous arsenic capping layer was deposited. The samples were then
transferred to STM and coupled BEEM experimental setup, where the amorphous As capping was
removed in situ at 435 °C. STM Tips were cleaned in situ by thermal heating (see ref. [5] for details).
Figure 1(a) shows atomically-resolved STM images obtained on a single quantum dot
surrounded by the wetting layer (WL). A clear (113)B plane (1x2) reconstruction (As-rich) is observed
which is very similar to the results obtained in the InAs/GaAs system.[4]. Four well-defined facets are
01 1 1 1 1 1 1 planes which are also observed in
obtained corresponding to the ( 1 ), ( 0 ), ( 00 ) and ( )
InAs/GaAs in ref. [4]. However, additional facets are measured in QDs formed on InP(113)B which
are not present in the InAs/GaAs system and not detected in previous TEM studies.[6] The welldefined
shape of QDs basis shows the complete absence of any rounded structure, as observed in ref.
[4], and attributed to different facet growth kinetics. We assumed that to be due to the lower lattice
mismatch in the InAs/InP system, and to long GI under As.
Finally, we investigated the electronic properties of individual InAs quantum dots by BEEM.
For this purpose, an Au Schottky contact was formed on the QDs layer. The buried nanostructures
were observed in the BEEM imaging mode. The energy of conduction band minima (, X and L
valleys), determined in the spectroscopy mode [7] are discussed
We acknowledge C’ nano Nord Ouest, Région Bretagne and Rennes Métropole for fundings.
[1] D. Bimberg, J. Phys. D: Appl. Phys., 38, 2055 (2005).
[2] T. J. Krzyzewski and T. S. Jones, Phys. Rev. B, 78, 155307 (2008).
[3] C. Cornet, M. Hayne, P. Caroff et al., Phys. Rev. B, 74, 245315 (2006).
[4] Y. Temko, T. Suzuki, P. Kratzer and K. Jacobi, Phys. Rev. B, 68, 165310 (2003).
[5] S. Guézo, P. Turban, C. Lallaizon et al., Appl. Phys. Lett., 93, 172116 (2008).
[6] D. Lacombe, A. Ponchet, S. Fréchengues et al., Appl. Phys. Lett., 74, 1680 (1999).
[7] S. Guézo, P. Turban, S. Di Matteo et al., Phys. Rev. B, 81, 085319 (2010).
__________________________
* Contact: charles.cornet@insa-rennes.fr

MoP26
Fig 1: (a) 100*76 nm² STM image (constant current mode of operation) of a single quantum dot showing a typical
As-rich (1*2) surface reconstruction of the WL. (b) 60*60 nm² STM image (constant current mode of operation) with
crystallographic facets. Absence of rounded structure is observed due to low lattice mismatch and GI.

MoP27
Investigations of growth kinetics of InN using pulsed RF MBE
A. Kraus * , R. E. Buß, H. Bremers, U. Rossow, and A. Hangleiter
Institute of Applied Physics, Technische Universität Braunschweig, Germany
For applications such as solar cells InGaN layers with indium concentrations of 10% to 100% are
needed which cover the near infrared to visible region. In a previous paper we have shown that
we can grow by RF-MBE InGaN layers with In concentrations up to 26% of high quality.
Here we concentrate on InGaN layers with high concentrations and the binary InN. Growing such
layers especially on high quality GaN template layers is difficult due to the high strain and
necessary balancing of the fluxes to avoid the segregation of In. Structural and morphological
defects may form as a sonsequence of strain relaxation. Furthermore, group-III-nitride layers
usually benefit from high growth temperatures well above 800°C, which in the case of InN
cannot be applied due to the relative weak In-N bond and consequently low nitrogen desorption
temperature.
In this contribution we report on the continued and indium pulsed mode growth of InN on GaN
templates in a Riber 32 P RF MBE reactor. Pulsing the growth was carried out by open the In
shutter only for a few seconds followed by a period where only active nitrogen reaches the
surface. For GaN, AlN and InGaN this is an effective method to grow high quality material [1,2].
This method is comparable with the MEE method, but there the nitrogen flux is pulsed too [3].
The conditions like pulse length and frequency, In flux, nitrogen flux and growth temperature
were varied to find optimal growth conditions. The growth was monitored in-situ by 633nm
reflectometry applied at an angle of incidence around 75° and by conventional RHEED.
During the time the In shutter is open we observe an increasing intensity of the reflected light.
This is depicted as part A in figure 1. Directly after closing the shutter the slope of this increase
becomes smaller or even saturates (part B in figure 1). This is followed by a decrease of the
intensity until the In shutter is opened again and this sequence restarts (part C).
Figure 1: Intensity variation of the reflected Laser beam. Figure 2: ω-2θ scans of two InN samples grown with
continious In flux (black) an pulsed In flux (red).
The samples were characterized by HR-XRD and AFM. XRD ω - scans showed an improvement
of the layer structure in terms of the FWHM. Compared to layers grown with a continuous supply
of In atoms the FWHM is approximately 4 times higher than for the samples grown with a pulsed
In flux. Furthermore the samples grown with a pulsed In flux are fully relaxed whereas samples
grown with a continuous In flux are significantly strained. AFM investigations of the surface
morphology exhibit smoother surfaces compared to the samples grown with a continuous flux.
This can be seen as proof for the improved quality of the InN layer.
[1] M. Moseley, J. Lowder, D. Billingsley, and W. A. Doolittle, Appl. Phys. Lett., 97, 191902 (2010).
[2] G. Namkoong et al., Appl. Phys. Lett., 93, 172112 (2008).
[3] H. Lu, W. J. Schaff, J. Hwang, H. Wu, W. Yeo, A. Pharkya, and L. F. Eastman, Appl. Phys. Lett., 77, 2548 (2000).
________________________________________
* Contact: a.kraus@tu-bs.de

MoP28
Low Thermal Budget Fabrication of Local Artificial Substrates by
Droplet Epitaxy on Silicon
S.Bietti, C.Somaschini, N.Koguchi and S.Sanguinetti
L-NESS and Dipartimento di Scienza dei Materiali, Università di Milano Bicocca, via Cozzi 53, 20153 Milano,
Italy
The droplet epitaxy (DE) growth method for the fabrication of III-V material quantum nanostructures
[1], is an intrinsically Low Thermal Budget technique, being fully performed at temperature between
200 and 350 °C. This makes DE perfectly suited for the realization of growth procedures compatible
with back-end integration. In short, the DE growth procedure consists in the deposition at different
times for the group III and group V elements. Group III elements create a regular pattern of liquid
droplet on a substrate, group V elements are incorporated inside group III element crystalling the
droplets into a quantum nanostructure.
We can distinguish two main areas where fabrication of III-V quantum nanostructures on Si substrate
could play a fundamental role. The first is the fabrication of nanostructured active layers at LTB with
designed DOS for optimum device performance [2]. The second area concerns the realization of local
artificial substrates for heterogeneous integration of quantum nanostructures [3].
(a)
(b)
Fig 1: density ans size of GaAs quantum nanostructures grown by DE on Si substrate in different conditions (a) and
SEM image of islands on one of the samples (b).
The nucleation of quantum dots atop an island is an attractive approach to address radiative
recombination issues and dot uniformity as the island both separates the dot from the interface with the
substrate and provides a nucleation platform of sufficiently small dimension to realize quantum size
effects. For this purpose we fabricated self-assembly of GaAs islands by DE which show highly
tunable density (from 10 7 to 10 9 cm -2 ) and size (from 75 nm to 250 nm) and size dispersion below
10%. Changing the substrate temperature during the Ga deposition and the amount of irradiated Ga is
possible to independently control the density and the size of the nanostructures (figure 1a). The
islands, made by single relaxed crystals, show well defined shapes, with a high aspect ratio (Figure
1b). The low thermal budget required for the island self-assembly, together with the high scalability of
the process, make these islands good candidates for local artificial substrates on Si.
[1] N. Koguchi and K. Ishige, Japanese Journal Of Applied Physics 32, 2052-2058 (1993).
[2] S.Bietti, C.Somaschini, S. Sanguinetti, N. Koguchi, G. Isella, and D. Chrastina, Applied Physics Letters 95, 241102
(2009).
[3] C. Somaschini, S. Bietti, N. Koguchi, F. Montalenti, C. Frigeri, and S. Sanguinetti, Applied Physics Letters 97, 053101
(2010).
__________________________
* Contact: sergio.bietti@mater.unimib.it

Integrated PL intensity (arb. units)
MoP29
Investigation of improvement and degradation of thermal
annealed indium rich InGaN/GaN quantum wells grown by NH 3 -MBE
N. A. K. Kaufmann 1,* , A. Dussaigne 1 , D. Martin 1 and N. Grandjean 1
1
Institute of Condensed Matter Physics, Ecole Polytechnique Fédérale de Lausanne (EPFL),
Station 3, CH-1015 Lausanne, Switzerland
While highly efficient blue InGaN/GaN light emitting diodes (LEDs) or laser diodes (LDs) are
commercially available, LEDs and LDs emitting in the green wavelength range are still challenging to
fabricate. Apart from the difficulty to increase the indium content above 25-30%, while maintaining a
good crystalline quality, degradation of high indium content quantum well (QW) layer may occur
during p-type layer deposition by metal organic vapor phase epitaxy (MOVPE) due to high growth
temperature (above 1000°C) [1].
Molecular beam epitaxy (MBE) allows realizing p-type layers at much lower temperatures (

MoP30
Shape Changes in Patterned Planar InAs as a Function of
Thickness and Temperature
K. G. Eyink 1,* , L. Grazulis 1 , K. Mahalingham 1 , M. Twyman 1 , J. Shoaf 1 , V. Hart 1 ,
J. Hoelscher 1 , C. Claflin 1 , and D. Tomich 1
1
Air Force Research Laboratory, AFRL/ RXPS, 3005 Hobson Way, Wright Patterson AFB, OH, USA
Quantum dots have the potential to produce devices with enhanced properties. However, many
quantum dot devices require the quantum dots to have a precise size and a precise location for
optimum operation. So far approaches such as directed assembly and self assembly have failed due to
the random effects resulting during nucleation of the quantum dots. InAs grown under metal rich
conditions can remain planar as opposed to forming the self assembled quantum dot morphology.
Recently we have demonstrated that planar InAs when patterned via tip-based scribing and then
annealed under an As pressure typical for self-assembled quantum dot growth reorganizes and assumes
a 3D morphology. We have been studying this process as a potential method to precisely locate
quantum dots with definable sizes. In this work we report change in the morphology for different
thickness of planar InAs for various pattern dimensions and annealing temperatures. We have analyzed
the composition of the films after annealing to determine the effect induced in the films from
patterning resulting from scribing. Using this approach, arrays of 3D InAs mounds have been formed
with mounds having a base dimensions of 800, 500, and 350Å. These results demonstrate that the
smaller patterns are less stable and coarsening becomes more dominant.
Many quantum dot devices require quantum dots to have precise size and location for optimum
operation. So far approaches such as directed assembly and self assembly have failed due to random
effects resulting during nucleation of the QDs. We have demonstrated that planar InAs when
patterned via tip-based scribing and then annealed under an As reorganizes and assumes a 3D
morphology. In this work we report changes in morphology for different thickness of InAs for various
pattern dimensions and annealing temperatures. Arrays of 3D InAs mounds have been formed with
mounds having a base dimensions of 800, 500, and 350Å.

MoP30
Occurrences
160
140
120
100
80
60
40
20
(a)
Area
0
0 500 1000 1500 2000 2500 3000 3500
Area (nm 2 )
(c)
75 line grid
Occurrences
160
140
120
100
80
60
40
20
(b)
Area
0
0 500 1000 1500 2000 2500 3000 3500
Area (nm 2 )
(d)
100 line grid
Occurrences
160
140
120
100
80
60
40
20
Height 75 line grid
0
0 5 10 15 20 25
Height (nm)
Occurences
160
140
120
100
80
60
40
20
Height 100 line grid
0
0 5 10 15 20 25
Height (nm)
(e)
(f)
Figure 1. 2μm x 2μm AFM images of annealed planar InAs patterns produced using (a) 75
and (b) 100 lines per pattern. The

MoP31
Influence of Al on the group III-assisted growth of axial
AlGaAs/GaAs heterostructure nanowires
T. Rieger * , M. I. Lepsa , H. Lüth , T. Schäpers and D. Grützmacher
Institute of Bio- and Nanosystems (IBN-1) and JARA-Fundamentals of Future Information Technology,
Forschungszentrum Jülich, 52425 Jülich, Germany
Nanowires (NWs) are promising candidates for future (opto-) electronic devices especially, when they
include controlled heterostructures. The self-catalyzed as well as Au-catalyzed growth of core-shell
GaAs/AlGaAs nanowires has been reported and axial heterostructure nanowires with the change of
group V element have been demonstrated as well. But, there are only a few reports about axial
heterostructure nanowires with different group III elements, mainly focusing on the GaAs/InAs system
[1] and/or Au-catalyzed growth [2].
In this work we report on the self-catalyzed growth of axial AlGaAs/GaAs heterostructure NWs on
HSQ (Hydrogen Silesquioxan) spin-coated GaAs (111)B substrates. We have investigated the
influence of Al growth rate, growth time and substrate temperature on the growth of axial
AlGaAs/GaAs heterostructure NWs. It is found that even small amounts of Al reduce the axial growth
but strongly promote growth on the amorphous oxide and NW sidewalls leading to unintentionally
grown core/shell NWs as it can be seen in the scanning electron micrograph in Fig. 1(a) for an Al
amount corresponding to Al 0.12 Ga 0.88 As.
The NW growth rate is influenced by the amount of Al. In Fig. 1 (b) the length of the NWs after 45
min of GaAs growth followed by 45 min of AlGaAs growth is plotted as a function of the Al amount.
For Al 0.12 Ga 0.88 As, the NW growth rate is reduced to 0.2 nm/s compared to 0.5 nm/s for GaAs, thus by
a factor of three. In fact, this is still 7 times the layer growth rate and demonstrates the possibility to
grow axial AlGaAs/GaAs heterostructure nanowires using self-catalyzed growth. For Al 0.21 Ga 0.79 As
the axial growth rate is close to zero. Also the switching back from AlGaAs to GaAs growth is found
to be difficult mainly due to the growth on the oxide and thus, reduced growth selectivity.
At growth temperatures of 670°C , above the Al melting point, the growth of pure AlAs NWs is not
possible using the self-catalyzed growth method.
(a)
(b)
Fig 1: (a) GaAs/AlGaAs NW with enhanced growth on the oxide, (b) length of AlGaAs/GaAs
heterostructure NWs after 45 minutes of GaAs and 45 minutes of AlGaAs growth
[1] M. Heiß et al., Nanotechnology, 20, 075603 (2009).
[2] M. Paladugu et al., Small, 3, 1873 (2007).
__________________________
* Contact: t.rieger@fz-juelich.de

MoP32
Ferromagnetic and transport properties of very thin (Ga,Mn)As
layers
L. Ebel * , F. Greullet, T. Naydenova, J. Constantino, S. Mark, C. Gould, K. Brunner
and L.W. Molenkamp
Physikalisches Institut (EP3), Universität Würzburg, Am Hubland, 97074 Würzburg, Germany
The strongly anisotropic valence band structure of the diluted magnetic semiconductor (Ga,Mn)As is
strongly influenced by the biaxial strain and hole density. By controlling one of these parameters it is
possible to change the magnetic anisotropies, which was successfully shown e.g. by uniaxial strain
relaxation of nanopatterned structures [1,2]. Another possibility to change the magnetic anisotropies
reversibly might be electrical gating a thin (Ga,Mn)As layer to control the hole density.
In thin layers (≤ 10nm), surface defects and point defects play an important role for the ferromagnetic
properties and lead to hole compensation, valence band bending and therefore to a reduced hole
density. In our III-V UHV-MBE chamber we have grown thin (~5nm) layers with homogeneous or
parabolically graded Mn content by Mn flux sequences with an accuracy of tenth of seconds during
the growth to obtain reproducibly conductive, ferromagnetic layers. This technique offers a central
(Ga,Mn)As layer of 5 ML and three Mn doping-spikes layers above and underneath (Fig. 1 (a) inset).
These spikes are used to shield the center part from the aforementioned defects. The samples were
investigated by RHEED, SQUID, room- and low-temperature 4-terminal transport measurements, as
well as simplified self consistent band alignment model calculations by “nextnano”.
The results show a strong influence of hole compensation and band bending caused by adjacent LT-
GaAs and surface defects on hole density. This calculated parabolic valence band structure for such
thin layers promises to be useful for gateable structures (Fig. 1 (a)). Thin layers with graded Mn
content reveal reproducibly better conductivity and higher Curie temperatures (Fig 1 (b)) compared to
homogeneous thin layers. These layers were grown under our standard growth conditions
(BEP(As 4 /Ga)=25, T sub =270°C) with nominally the same Mn content (4%).
Using these results as a starting point, we have optimized the growth conditions for thin parabolic
graded (Ga,Mn)As layers with lower Mn content of x=2,5% by varying BEP(As 4 /Ga) and T sub (Fig. 1
(c)) to reduce the As-Antisite and/or Mn interstitial concentration which are compensating deep
donors in low-temperature (Ga,Mn)As growth. We found a distinct growth window at lower As-flux
(BEP(As 4
/Ga)=20) and higher substrate temperature (T sub
=290°C) compared to thick standard layers,
in which it is possible to grow thin layers with well-controlled transport and ferromagnetic behavior.
The conductivity and Curie-temperature (T=41K) of these as-grown layers are comparable to thicker
reference (Ga,Mn)As layers. We conclude that the growth conditions and the layer structure for such
thin layers play a crucial role.
[1] K.Pappert, S.Hümpfner, C.Gould, J.Wenisch, K.Brunner, G.Schmidt and L.W.Molenkamp, Nature Physics, 3, 573-578
(2007).
[2] J.Wenisch, C.Gould, L.Ebel, J.Storz, K. Pappert, M.J. Schmidt, C.Kumpf, G.Schmidt, K.Brunner and L.W.Molenkamp
PRL, 99, 077201 (2007).
__________________________
* Contact: lebel@physik.uni-wuerzburg.de

MoP32
Fig 1 (a): Self consistent solution of Poisson- and Schrödinger-equation for a parabolically graded Mn content sample. The
shape of the valence band edge is roughly parabolic. (inset): layer structure
Fig 1 (b):.Comparison of the Curie temperature of a bulk (70nm), homogeneous (4nm) and parabolically graded Mn content
layer (4nm of effective thickness of (Ga,Mn)As). All samples have a nominally Mn concentration of 4%.
Fig 1 (c): Low temperature conductivity (measured in 4 terminal Hall geometry) of parabolically graded (Ga,Mn)As layers
with a Mn content of 2.5 %, as a function of the substrate temperature T sub
and the BEP ratio (As 4
/Ga).
Optimal growth, i.e. highest conductivity and a Curie-temperature of 41K (not shown), occurs around 290°C and at a
BEP ratio (As 4 /Ga)=20. It results in good, well-controlled, electrical and ferromagnetic properties.

Self-assembled InP-nanoneedles grown on (001) InP by gas source
MBE
M. Chashnikova 1 , V. Bryksa 2 , A. Mogilatenko 1 , O. Fedosenko 1 , S. Machulik 1 , M.P.
Semtsiv 1 , W. Neumann 1 , and W.T. Masselink 1
1
Department of Physics, Humboldt University Berlin, Newtonstr. 15, 12489, Berlin, Germany
2 Insitute of Semiconductor Physics, NASU, Nauki pr. 41, 03028 Kiew, Ukraine
MoP33
One dimensional semiconductor crystals attract a great deal of attention due to both, the prospect of
new physical effects and the potential for application in nanoelectronic devices [1-5]. These structures
have unique optical and electronic properties stemming from their one-dimensional crystalline
structure and large surface-to-volume ratio. Nanoneedles (NNs) are structures characterized through
their sharp tip and narrow taper that may have advantages for potential applications as transistors [6-9]
and in nonlinear optics such as tip-enhanced Raman spectroscopy [10]. Such sharp tips are apparently
only possible to achieve through catalyst-free growth. GaAs NNs were recently grown by MOVPE
[11]. This work presents InP-NNs grown catalyst-free by gas source molecular beam epitaxy
(GSMBE) on InP (001) substrates. Harvested NNs were investigated by transmission electron
microscopy (TEM), Raman spectroscopy (RS), photoluminescence (PL) and cathodoluminescence
(CL) spectroscopy to obtain valuable information about their bulk and surface microstructure.
Self-assembled NNs were grown in a Riber Compact 21T GSMBE system equipped with arsine and
phosphine as group-V elements sources. High purity (6N5 and 7N) solid In and Al were used as the
group-III element sources. Before growing InP NNs a 1µm-thick InAlAs layer was grown on the (001)
InP substrate and then exposed to air at ambient pressure for several days to produce a thin top layer of
native oxides. The sample was then heated at 480°C in the GSMBE chamber under phosphine flux
allowing for desorption of In, As and P oxides. Al oxide is likely to withstand this thermal treatment
and will mask part of the surface. Subsequent growth of InP localized at the oxide-free spots started at
460°C at a very low rate of 0.01 nm/s, then, steadily increased to 0.3 nm/s. InP layers were Si-doped to
5×10 18 cm -3 , and the total amount corresponds to a 1µm thick layer in a two dimensional case.
Scanning electron microscopy (SEM) revealed that the above mentioned amount of InP developed into
densely packed NNs (presumably of hexagonal cross section) with an average base of 1000 nm and a
length of 30-40 µm (see Fig. 1a). Two types of NNs are observed: a) NNs tilted to the surface at
approximately 54.74° (likely grown along direction) and b) NNs vertically oriented (see Fig.
1a). Tilted NNs outnumber the [001]-oriented ones by an average ratio of about 4:1. The vertical and
the tilted NNs can be distinguished from each other even when removed from the substrate, since only
the tilted ones show a faceted shape of the surface directly exposed to In and Ga fluxes (see Fig. 1a).
The crystalline structure of the tilted NNs, studied by TEM, reveal a mixture of ZB and WZ regions
(Fig. 1b). The hexagonal WZ InP regions are larger than the cubic ZB InP segments. We assume that
the surface faceting of the tilted NNs originate from the mixed − ZB and WZ − structures. So the rare
NNs grown along [001] axis have likely a pure ZB structure, from Raman measurements (see figure
2), as they do not show any faceting on the side walls.
[1] M.T. Björk, B.J. Ohlsson, C. Thelander, A.I. Persson, K. Deppert, L.R. Wallenberg, L. Samuelson, Appl. Phys. Lett. 81
4458 (2002)
[2] X. Duan, Y.Huang, R. Agarwal, C. Lieber, Nature 421 (2003) 241
[3] M. Yazawa, M.Koguchi, A.Muto, K.Hiruma, Advanced Materials 5 (1993) No.7/8
[4] C. Thelander, T. Martensson, M.T.Björk, B.J. Ohlsson, M.W. Larsson, L.R. Wallenberg, L. Samuelson, Appl. Phys.
Lett.2002, 81, 4458
[5] Y.Huang, X.Duan, Y.Cui, and C.M.Lieber, Nano Letters 2(2002)101
[6] Tans, S. J.; Verschueren, R. M.; Dekker, C. Nature 1998, 393, 49
[7] Martel, R.; Schmidt, T.; Shea, H. R.; Hertel, T.; Avouris, P.Appl. Phys. Lett. 1998, 73, 2447
[8] Zhou, C.; Kong, J.; Dai, H. Appl. Phys. Lett. 2000, 76, 1597
[9] C.P.Collins, M.S.Arnold, P. Avouris, Science 2001, 292, 706
[10] R.M. Stöckle, Y.D.Suh,V.Deckert, R.Zenobi Chem.Phys.Lett. 318 (2000) 131
[11] M. Moewe, L.C. Chuang, S. Crankshaw, C. Chase and Connie Chang-Hasnain, Appl. Phys. Lett. 93 (2008) 023116

MoP33
Fig.1 (a) SEM-image of extra long InP-NNs grown on InP (001) substrate (b) High resolution TEM image of a
tilted heterostructured InP-NN
Fig.2 Micro-Raman spectra obtained at InP nanoneedles: Spectra (1) and (2) were acquired at different positions
along an inclined nanoneedle consisting of an alternating ZB/WZ structure. Spectrum (3) indicates the presence
of the nanoneedle with a pure ZB structure.

MoP34
Selective growth of InP on pre-patterned wafers by the means of
Gas-Source MBE
A.Aleksandrova 1 , G.Monastyrskyi 1 , O.Fedosenko 1 , M.Chashnikova 1 , S.Machulik 1 ,
J.Kishkat 1 , M.P.Semtsiv 1 and T.W.Masselink 1
1
Humboldt University Berlin, Department of Physics, Newtonstr.15,12489 Berlin, Germany
Keywords: Growth from the high temperature solutions, infrared devices, semiconducting InP,
semiconducting Al compounds, MBE, pre-patterned growth.
Gas-source Molecular Beam Epitaxy (GSMBE) has been used to grow InP selectively on InP:S
substrates pre-patterned with SiO 2 masks. In this paper we discuss the wafer preparation process and
growth temperature dependence on the sticking of InP on SiO 2 mask and morphology of the grown
layers, as well as potential application in optoelectronics.
The SiO 2 mask has been prepared on epi-ready InP:S substrates in three following steps: (i) formation
of 40µm-wide photo-lack stripes by the means of conventional optical lithography, (ii) reactive
magnetron sputtering of SiO 2 film, and (iii) conventional lift-off process and the wafer cleaning. The
wafer cleaning process has involved (a) remover rinsing, (b) developer rinsing, (c) plasma oxidation
and optionally (d) cleaning with either HF or H 2 SO 4 :H 2 O 1:1 solution.
Growth of InP test samples has been carried out between 450°C and 550°C measured by
thermocouple. Selective growth of InP on InP (in the SiO 2 windows) has been observed at the
temperature of 500°C and above. However growth temperatures much higher then 500°C has led to
rough surfaces, likely due to In-stabilized growth front (Fig. 1). Laser ridge, defined by wet chemical
etching, has been overgrown at 500°C with a smooth surface and without visible interface defects.
Further in the paper we will discuss applicability of GSMBE technique for development of buriedheterostructure
lasers.
__________________________
* Contact: anna_a@physik.hu-berlin.de
Fig 1. SEM image of InP grown on InP in the stripe-windows of SiO2 mask at 520°C.

MoP35
Scaling of quantum cascade laser efficiency
with a number of cascades
O.Fedosenko 1 , A.Aleksandrova 1 , G.Monastyrskyi 1 , M.Chashnikova 1 , S.Machulik 1 ,
J.Kishkat 1 , M.Klinkmüller 1 , M.P.Semtsiv 1 and T.W.Masselink 1
1
Humboldt University Berlin, Department of Physics, Newtonstr.15,12489 Berlin, Germany
Keywords: Growth from the high temperature solutions, infrared devices, semiconducting InP,
semiconducting Al compounds, MBE.
Quantum cascade lasers (QCLs) [1,2,3] has become the ultimate compact mid-infrared laser solutions
for a broad variety of applications, dominated mostly by gas spectroscopy. In this paper we focus on
scaling of the QCL slope efficiency, threshold current density, and other relevant parameters with
increasing the number of cascades.
The most interesting for us in this case is QCL with high cascade number. The way to increase output
optical power there will be to increase the volume of the laser active region, the doping or simply the
input electrical power. In this paper we described a case with volume increasing, namely laser core
thickness (higher cascade number).
QCL samples were grown in a Compact 21T gas-source MBE system from Riber. The system is
equipped with high-temperature cracking cell, providing radicals of AsH 3 and PH 3 as a source of
group-V elements. Group-III elements are supplied by conventional effusion cells equipped with 80cc
conical crucibles. The crucibles were loaded with approximately 10gr., 30gr., and 30gr. of Al, In, and
Ga respectively. The growth sequence starts with low doped (n=1-2·10 17 cm -3 ) InP:S substrate and
follows with InP buffer, InGaAs spacer, InGaAs-InAlAs multilayer active zone, InGaAs upper spacer,
upper InP optical cladding, and ends up with highly doped InGaAs contact layer. Detailed recipe is
given by Green et. al [4]. We have grown a series of samples with different number, N casc. , of emitting
cascades (periods) in active zone. Growth rate of all used materials was kept between 0.2 and 0.3
nm/s. N casc. has ranged between 25 and 200, which has result in the growth time of the active zone
between 2 and 15 hours. Further in the paper we will discuss dependence of electrical parameters on
cascade number (Fig.1).
Slope efficiency (W/A)
2
1
0
0 50 100 150 200
Cascade number
Fig 1: (a) Slope efficiency as a function of cascade number, RT
[1] A. F. Kazarinov and R. A. Suris, Sov. Phys. Semicond. 5, 207 (1971).
[2] J.Faist, F.Capasso, D.L.Sivco, C.Sirtori A.L.Hutchinson, and A.Y.Cho, Science 264, 553 (1994).
[3] C. Gmachl, F. Capasso, D. L. Sivco, A. Y. Cho, Rep. Prog. Phys. 64, 1533 (2001).
[4] R. P. Green et al, Appl. Phys. Lett. 85, 5529 (2004).
__________________________
* Contact: oliana@physik.hu-berlin.de

MoP36
MBE growth of quantum-cascade laser on pre-patterned
substrates
G. Monastyrskyi, O. Fedosenko, M. Chashnikova, A. Alexandrova, S. Machulik,
J.Kischkat, M. Klinkmüller, M. P. Semtsiv and W. T. Masselink
1
Department of Physics, Humboldt University Berlin
Newtonstr. 15, 12489 Berlin, Germany
Keywords: quantum cascade laser, MBE, pre-patterned substrate
In this paper we present our results on growth InP and lattice matched InGaAs-InAlAs on prepatterned
substrates. The idea is to grow together monocrystalline and polycrystalline structure on the
same substrate. One of ways to do it is to cover part of substrate by some amorphous material. In our
case was used SiO 2 stripes 50 µm wide with periodicity 500 µm, which was inflicted on InP substrate.
Thus prepared InP substrate was used to grow semiconductor structures in gas-source MBE. For this
samples desorption was observed at temperatures about 50° C higher then usual for InP substrates in
our system.
Series of samples was grown at different substrate temperature to determine optimal conditions of
growth. Figure 1a,b shows InP-InGaAs-InAlAs test structures cleaved across the SiO 2 stripes.
(a)
(b)
Fig 1: InP-InGaAs-InAlAs structures, (a) grown at 560°C, (b) grown at 475°C
Sample on Fig. 1(a) is grown at 560°C for InP layer and 490° C for InGaAs-InAlAs layer; sample on
Fig. 1(b) is grown at 475°C for InP layer and 400° C for InGaAs-InAlAs layer. Growth at 560°C
provides pretty smooth surface of the crystalline material and does not reveal any interface defects.
This growth temperature and the substrate preparation procedure including cleaning with diluted
H 2 SO 4 were used for the growth of quantum-cascade structure on pre-patterned InP:S substrate.
Approx. 35-µm wide laser stripes were processed above the windows in the SiO 2 mask. Processed
QCLs operate at ~11µm wavelength at room temperature. Threshold current density at room
temperature is ~ 3kA/cm 2 and at 80K is ~ 1kA/cm 2
This QCL fabrication shows a feasibility of QCL fusion with integrated optical circuits.
__________________________
* Contact: grygorii@physik.hu-berlin.de

MoP37
Non polar GaN/ZnO heterostructures grown by ammonia source
molecular beam epitaxy
J. Brault 1,* , G. Sophia 1 , S. El Kazzi 1 , J.-M. Chauveau 1,2 , P. Vennéguès 1 , M. Nemoz 1 ,
M. Teisseire 1 , M. Leroux 1 , C. Deparis 1 , C. Morhain 1 , O. Tottereau 1 , L. Nguyen 1
1
CRHEA-CNRS, Rue Bernard Grégory, 06560 Valbonne, France
2 Phys. Dept., University of Nice Sophia-Antipolis, 06103 Nice, France
Although heteroepitaxially grown on sapphire, gallium nitride (GaN) based optoelectronic devices
have reached outstanding performances in the near UV – blue range: for instance, state-of-the-art blue
light emitting diodes (LEDs) have internal quantum efficiencies (IQEs) around 90%. However, these
LEDs are fabricated with heterostructures grown on the polar (0001) “c-plane” orientation which leads
to the presence of an internal electric field in their active region (typically several hundreds of kV/cm).
These large electric fields can have a strong impact on LED characteristics (quantum confined Stark
effect), and as a consequence, it still remains a challenge to realize LEDs with large IQEs in the UV
and green-yellow range. Therefore, the use of nonpolar orientations, allowing for a suppression of the
internal electric field along the growth direction [1], have recently attracted much attention for device
fabrication on nonpolar bulk GaN substrates [2].
However, the growth of nonpolar GaN is currently facing two major difficulties: i) the large densities
of defects in GaN nonpolar layers grown on host substrates (sapphire, SiC…), and ii) the scarcity and
high cost of nonpolar GaN substrates for homoepitaxy. Another approach could be the use of nonpolar
zinc oxide (ZnO) substrates, since this material has the same crystallographic structure as GaN, close
lattice parameters and similar thermal expansion coefficients. Moreover, nonpolar high quality ZnO
substrates are commercially available and affordable, although the size is still limited to a few cm 2 .
However, the thermal stability of ZnO could be a limiting factor for its use with nitride materials,
which usually require growth temperatures above 1000°C. Our objective is then to develop a low
temperature growth process for the realization of nonpolar GaN/ZnO heterostructures by using
molecular beam epitaxy (MBE). As a first step, in this study, we have grown GaN layers on nonpolar
(11-20) “a-plane” ZnO/sapphire templates. These templates have been fabricated in CRHEA by MBE
[3]. At first, we have determined, by X-Ray Diffraction, that the GaN layers have the expected (11-20)
orientation (fig. 1(a)), and present a similar morphology as the ZnO templates [3], with stripes
elongated in the [0001] direction (fig. 1(b)). Transmission electron microscopy has also been
performed and we have observed that the structural quality of GaN is fairly related to the structural
characteristics of ZnO. We have also studied the influence of growth parameters on the structural and
optical properties of GaN: for instance, depending on the growth temperature, GaN band-edge
emission can be clearly observed or strongly reduced, as presented on fig. 2. Based on these results,
the optimization of the growth conditions for the MBE growth of GaN on ZnO will be discussed.
Intensity (a.u.)
10 5 Sapphire S -204
-102
(a)
S -408
GaN 110
ZnO 110
GaN 220
10 4
S -306
10 3
10 2
10 1
ZnO 220
10 0
20 40 60 80 100 120 140
2θ/ω scan (°)
(b)
2 µm
Intensity (u.a.)
10 GaN band-edge
5
transition
10 4
Donor-Acceptor
pairs transitions
Basal Stacking
Faults emission
10 3
10 2
λ exc.
= 244 nm
P exc.
= 25 mW
700°C
650°C 600°C
T = 12K
350 360 370 380 390
Wavelength (nm)
Fig 1 (left-side): (a) X-Ray 2θ/ω scan from a GaN layer grown on ZnO/sapphire, (b) AFM image of a GaN layer.
Fig 2 (right-side): Photoluminescence spectra from a-plane GaN layers grown at different temperatures.
[1] for instance Z. Bougrioua et al., Phys. Stat. Sol. (a), 204, 282 (2007).
[2] for instance M. C. Schmidt et al., JJAP Part 2-Letters & Express Letters, 46, L190 (2007).
[3] J.-M. Chauveau et al., J. Crystal Growth, 301-302, 366 (2007).
__________________________
* Contact: jb@crhea.cnrs.fr

MoP38
Emission of colloidal nano-crystals embedded in MBE grown ZnSe
microstructures
J. Kampmeier 1 , M. Rashad 1, 2 , A. Pawlis 1 , D. Schikora 1 , K. Lischka 1,*
1 Department Physik, Universität Paderborn, Warburger Str. 100, 33098 Paderborn, Germany
2 Department of physics, University of Assiut, 71516 Assiut, Egypt
Colloidal Nano-Crystals (NCs) exhibit a great potential regarding the tuning of their optical and
structural properties. The size, shape, material composition and the density of the NC can be varied
and therefore exactly matched to the requirements of specific optoelectronic devices. Alternatively,
different types of NC which emit light at several wavelengths can be combined in the same device to
achieve multi-color or white light emission from a single chip. In this context the key issue of all
applications based on colloidal NCs is the integration of the Quantumdots (QDs) in a matrix to
stabilize their optical properties and to facilitate the fabrication of micro-resonators with colloidal
NCs.
We have combined molecular beam epitaxy (MBE) of ZnSe with externally wet-chemically prepared,
colloidal NCs of CdSe to achieve fully integrated semiconductor structures grown by molecular beam
epitaxy [1]. This technique allows several degrees of freedom for choosing the NCs density, shape and
size. We used CdSe core, CdSe/ZnSe and CdSe/ZnS core/shell Nano-Crystals with shells of 1 to 2
mono-layers. The NCs were prepared in solution with a nominal concentration in volume of
10 12 NCs/l and 10 16 NCs/l, respectively. Using a spray-coating technique, the NCs were deposited on a
50 nm thick ZnSe buffer layer. Then, the sample was transferred back to the MBE chamber to
continue the growth of ZnSe by atomic layer epitaxy. Structural characterization by high resolution x-
ray diffraction revealed pseudomorphic ZnSe strained layers.
We find a distinct blue shift of the NC related photoluminescence (PL) when core/shell dots are
overgrown by ZnSe. This effect was explained by a model using the following assumptions: (i)
dissolution of the shell of the dots during the cap layer epitaxy; (ii) after overgrowth NCs are separated
from the ZnSe matrix by an interface barrier which prevents carriers generated in ZnSe to recombine
in the NCs.
We used Rapid Thermal Annealing (RTA) to remove the barrier between the NCs and ZnSe. The
effect of different annealing temperatures on the optical properties of CdSe core, CdSe/ZnSe and
CdSe/ZnS core/shell NCs overgrown with cap layer of ZnSe has been investigated. After RTA at
600 K the PL exhibits a red shift as compared to unprocessed samples. With the same model and the
same parameters applied for the overgrown NCs before annealing, the observed red-shift of the PL
after annealing can be explained by dissolution of the surface barrier between the NCs and the
surrounding ZnSe layer. For all three types of NCs we find an excellent quantitative agreement
between experimental and calculated data.
In order to study the PL of single or a small number of NCs we have structured the ZnSe layers using
e-beam lithography and wet chemical etching to produce micro-discs and micro-pillars which contain
a limited number of NCs. A micro-PL set up with a spatial resolution of about 1.5 µm was used. The
measured PL spectra clearly show the features of radiative recombination of excitons localized in
colloidal NCs.
[1] U. Woggon, et. al, Nano Lett. 5, 483 (2005).
__________________________
* Contact: lischka@upb.de

MoP39
Reproducible temperature calibration technique for GaAs MBE
François Morier-Genoud 1* and Denis Martin 2
1
Laboratoire d’OptoElectronique Quantique, Ecole Polytechnique Fédérale de Lausanne, Switzerland
2
Laboratoire en Semiconducteurs avancés pour la photonique et l’électronique, Ecole Polytechnique Fédérale
de Lausanne, Switzerland
In this work we present an improved technique to establish a temperature reference point for the GaAs
(100) substrate, based on changes in the RHEED reconstruction pattern from the Arsenic (As) to the
Gallium stabilized surfaces by varying the amount of As arriving on the surface.
We observe the transition occurring while decreasing the As species arriving on the surface at a
temperature higher than the congruent sublimation temperature. The As species at the surface, which
form a 2x4 reconstruction, evaporate leaving a 2x3 Ga rich reconstruction.
The transition time depends on the temperature of the substrate (Ts), the pressure of the As before
closing the valve, the leak rate of the closed As source and the As pumping speed.
A photo detector (camera) reports the disappearance of the signal of the half order streak of the 4 fold
2x4 reconstruction, while the beam incident angle is chosen to produce the maximized intensity from
the diffracted signal. When As source is turned off (shutter and valve closed), the reconstruction
changes as fast as As evaporates from the surface. This transition time is around one second for a
temperature of 590°C and increase to more then 10 seconds at 560°C.
The first graph shows the dependency of the transition time versus As pressure for three different
temperatures. As it could be seen for the highest temperature, the transition time stays constant for a
pressure above 1.2E-5 Torr. As it could be seen on the second graph, the reproducibility of this
method is reliable in a vide range of As pressure.
This rapid calibration method is not depending on the system configuration neither on surface
treatment of the substrate and could be helpful for all GaAs based MBE system.
Substrate temperature calibration
Fit for two As pressure ranges
10
10
Transition time [s]
Ts 588°C
Ts 580°C
Ts 575°C
1
2.5 10 -5 2.1 10 -5 1.7 10 -5 1.3 10 -5 9 10 -6 5 10 -6 1 10 -6
Transition time [s]
9
PAs= 8-9 E-6
8
PAs=1.2-2.1 E-5
7
6
5
4
3
2
1
0
560 565 570 575 580 585 590 595 600
Arsenic BEP [Torr]
__________________________
Substrate temperature [°C]
* Contact: francois.morier-genoud@epfl.ch

Tuesday Poster Session

TuP01
Fabrication and optical properties of CdTe quantum dots in ZnTe
nanowires
P. Wojnar 1,* , E. Janik 1 , A. Petroutchik 1 , L. Baczewski 1 , M. Goryca 2 ,
T. Kazimierczuk 2 , P. Kossacki 2 , G. Karczewski 1 and T. Wojtowicz 1
1
Institute of Physics, Polish Academy of Sciences, Al Lotników 32/46, 02-668 Warsaw, Poland
2 Institute of Experimental Physics, Warsaw University, ul HoŜa 69, 00-681 Warsaw, Poland
Quantum dots consisting of nanometer sized insertions of low energy gap semiconductor in large
energy gap nanowire have recently risen a great interest because of their emerging applications as
single photon sources at room temperature [1], biological markers, nanobarcodes [2] and also in view
of electronic coupling of several quantum dots. In parallel to this research, it has been shown that
CdTe quantum dots (grown by self assembly) containing a single Mn ion may act as a spin memory
unit with the possibility of optical writing and read out of the information on the spin of the Mn-ion
[3]. In this work, the fabrication and optical properties of CdTe quantum dots in ZnTe nanowires is
reported which is the first step towards coupling of several such spin memory units.
The structure is grown by molecular beam epitaxy by employing the vapor-liquid-solid growth
mechanism assisted by gold catalysts. Nanometer sized droplets of gold/gallium eutectic are formed
on GaAs substrate by thermal treatment of a 1 nm thick gold layer deposited in a separate growth
process. Subsequently, ZnTe nanowires with the length of about 1.5 µm are grown at relatively high
substrate temperature of 420°C. For the deposition of the CdTe insertion, the temperature is decreased
to 380°C, which is still a relatively high temperature, that ensures that CdTe growth occurs mostly in
the axial direction of the nanowire. Deposition time of CdTe is set to 60s. The further growth of
200 nm ZnTe nanowire segment takes place at 380°C.
Photoluminescence measurements performed at 5 K show clearly that the presence of CdTe insertions
in the ZnTe nanowires results in the appearance of a strong emission in the 2.0 eV – 2.2 eV energy
region. When the excitation spot diameter is reduced to the size of the order of 1 µm, this broad band
splits into several sharp lines with the spectral width of the order of 1 meV, which are related to the
emission from individual quantum dots in nanowires as confirmed by our further study. In order to
better resolve the individual emission lines, the nanowires are removed from the GaAs substrate in a
ultrasonic methanol bath and dropped onto a silicon substrate.
The sharp emission lines exhibit a relatively large degree of linear polarization ranging from 70% to
90% depending on the particular line. This effect is due to large contrast of the dielectric constant of
the nanowire and the surrounding and, thus, confirms that the emissions come from a nanowire
heterostructure. On the other hand, photon correlation measurements performed under continuous
wave excitation prove the zero dimensional character of studied objects. A clear antibunching at the
zero delay time between arriving of two subsequent photons from the same emission line is observed
confirming that these individual objects cannot emit simultaneously two photons at the same energy.
Moreover, the power dependence of the emission and the cross-correlation measurements, allow us to
indentify biexcitonic emissions with binding energies ranging from 7 - 12 meV
The research was partially supported by the European Union within European Regional Development Fund through grant
Innovative Economy (POIG..01.01.02-00-008/08)) and by the Foundation for Polish Science through subsidy 12/2007
[1] A. Tribu et al., NanoLett., 8, 4326 (2008)
[2] M.S. Gudiksen, L.J. Lauhon, J. Wang, D.C. Smith and C.M. Lieber, Nature 415, 617 (2002)
[3] M. Goryca et al. Phys. Rev. Lett. 103, 087401 (2009), C. Le Gall et al., Phys. Rev. Lett. 102, 127402 (2009)
__________________________
* Contact: wojnar@ifpan.edu.pl

TuP02
In-situ, real time Auger Monitoring as a new tool for growth
characterization and control
P. Staib
1
Staib Instruments Inc.,Williamsburg VA, USA
A new Auger Electron Spectrometer designed to monitor the surface composition during growth is
presented. The device is able to operate without impairing the deposition process and is compatible
with the environmental constraints of deposition chambers. The Auger electron excitation can be
generated by the electron beam used for RHEED studies thus allowing simultaneous elemental and
structural sample analyses. The Auger analyzer is mounted using one of the ports facing the sample,
preferably at normal incidence angle and the clearance between sample surface and analyzer is large
enough to clear the fluxes of material from multiple deposition source commonly used in growth
chambers. The analyzer has a variable energy resolution that can be electronically adjusted and can be
set for higher sensitivity and shorter acquisition time or better energy resolution. Examples of
applications are shown to illustrate the capabilities of the device.
Real time AES during ZnO growth on Sapphire and GaN substrates are monitored during growth by
AES and REELS. Figure 1 (curve a) shows the Auger spectra of Ga line from the substrate before
opening the Zn shutter. The Zn signal rapidly grows after opening the Zn shutter (curve b) and
remains fairly constant during growth. Oxygen Auger signals are shown in Figure 2. The O (KVV)
line is already observed after opening the Zn shutter and keeping the Oxygen RF source shutter closed.
The O2 residual gas pressure in the chamber is large enough to oxidize the deposited Zn (curve a).
The O (KVV) signal sharply increases after opening the Oxygen RF source shutter and reaches a peak
(curve b) at a sample temperature of 300 o C at which the RHEED shows a 3D like surface structure.
The O (KVV) signal decreases to a steady value once the sample temperature is raised to about 600 o C
and the surface becomes flat.
The Auger probe is used by Calley et al.[1] to the monitor the co-deposition of Fe, Dy, and Tb on
silicon substrates. The MNN Auger lines shown in Figure 3 correspond to pure Tb (curve a) and pure
Dy (curve b). In this example, these techniques are complicated by the fact that the two Auger lines
strongly overlap. The measured energy of an alloy Tb(70) Dy(30) is shown in curve c. It shows that
in spite of the overlap, there are energy regions, marked by the lines a and b, where the distributions
vary markedly between the pure elements. It is then possible to characterize a specific alloy
composition calibrating the signals in this energy range as shown by Calley et al. A sensitivity of 2%
for the variation of atomic ratios could be demonstrated.
In addition to AES, the analyzer provides REELS spectra of the distribution of characteristic energy
losses. Further capabilities of the analyzer probe, such as the measurement of the electron reflectivity,
are presented.
Acknowledgments
- Prof. H. Morkoc and Dr. V. Avutin for the testing during ZnO growth using an SVTA chamber.
- L. Calley, J. Lowder, W. Henderson and Prof. A. Doolittle for their help using the probe on a VEECO
GEN II chamber.
[1] L. Calley et al. Proceedings of the NAMBE2010 and this conference.
__________________________
*
Contact: pstaib@staibinstruments.com

TuP02
Figure 1 Figure 2
Figure 3

TuP03
Optimisation of Unusual Quantum Dot Growth Conditions
for Optical Coherence Tomography Applications
M. Hugues 1,2* , M. A. Majid 1 , S. Vezian 2 , D. T. D. Childs 1 , K. Kennedy 1 , and R. A.
Hogg 1
2 CRHEA-CNRS, rue Bernard Gregory, 06560 Valbonne, FRANCE
Finally the best QD SLDs have been used for in vivo OCT imaging of skin.
PL intensity (arb. units)
1.0
0.8
0.6
0.4
0.2
(a)
SLD
FWHM
=
120nm
Laser
FWHM
=
40nm
EL (dB)
(b)
0.1-1A pulsed
Max 3dB BW ~160 nm
0.0
1100 1200 1300 1400
Wavelength (nm)
1000 1100 1200 1300 1400
Wavelength(nm)
Figure 1. (a) Room-temperature PL spectra of QD structure optimized for laser (red) and SLD (blue) application.
(b) Emission spectrum for a QD SLD with the optimized active layer.
__________________________
* Contact: mh@crhea.cnrs.fr

TuP04
A New Route for Strain relaxation in In 0.52 Al 0.48 As on GaAs Grown
by MBE
Y. X. Song 1 , S. M. Wang 1* , Z. H. Lai 1 , M. Sadeghi 1 and J. R. Dong 2
1
Department of Microtechnology and Nanoscience, Chalmers University of Technology, SE-41296 Gothenburg,
Sweden. 2 Suzhou Institute of Nano-tech and Nano-bionics, CAS, Ruoshui Road 398, Suzhou 215125, P. R. China
Restricted by the availability of substrates with desired lattice constants, the design and epitaxy of most
semiconductor devices can only be based on nearly lattice matched material systems such as GaAs/AlAs
on GaAs substrate. Metamorphic growth is one of the solutions to overcome the lattice mismatch and
provides virtual substrates with arbitrary lattice constants and relatively low defect density. There are
large demands of In x Ga 1-x As/In x Al 1-x As metamorphic virtual substrates, for example, In x Ga 1-x As and
In x Al 1-x As buffers with about 10-15% or 25-30% In composition are suitable to fabricate 1.3 and 1.55
μm telecom lasers on GaAs with better performance than those on InP, while In 0.3 Ga 0.7 As is a candidate
material for 1 eV solar cells to replace the Ge cell (0.66 eV) to enhance efficiency of muti-junction solar
cells. However, growth of highly lattice mismatched (>2%) heterostructures usually leads to formation
of quantum dots and a high density of sessile edge-type threading dislocations. In this work we propose
a new growth scheme to grow relaxed In 0.52 Al 0.48 As on GaAs by MBE using low growth temperatures
combined with alternative supplies of cations and anions, resulting in smooth surface and strain
relaxation by 60° mixed dislocations. These 60° mixed dislocations are possible to be eliminated by
dislocation engineering methods.
Samples were grown at a growth temperature from 280 to 320 °C measured by a thermocouple. The first
40 ML In 0.52 Al 0.48 As was grown as follow. We first deposit 1ML mixture of In and Al without As to
obtain III-rich condition, and then open As for crystallization. The As is then closed followed by a 5s
pause before the new loop is started. It was observed from in situ reflection high-energy electron
diffraction (RHEED) that, for the first 25 loops, the streaky RHEED pattern became broad and short due
to formation of two dimensional (2D) islands
and then recovered slowly in the following loops.
No 3D islands were formed judged from
RHEED. After the 40 loops of InAlAs
deposition, a 500nm InAlAs bulk was grown in a
normal growth mode at 510°C measured by a
pyrometer. The final RHEED showed nice
streaky pattern indicating smooth surface. No
cross-hatch pattern, which is otherwise typical
for alloy graded buffers, was found from atomic
force microscopy (AFM) scans. It is observed
from transmission electron microscopy (TEM)
(Fig 1) that, the type of misfit dislocations (MDs)
is a mixture of both 60° and 90°. Most TDs are
60° type and few 90° type ones can be found.
This is promising for implementing TD blocking
schemes, which are normally effective only on
60° type TDs [1].
Fig 1. Cross-section TEM image of the
In 0.52 Al 0.48 As sample.
[1] Y. X. Song and S. M. Wang, et al. Appl. Phys. Lett. 97, 091903 (2010)
_________________________
* Contact: shumin@chalmers.se

TuP05
Nonpolar III-nitride microcavities for polariton lasing
A. Dussaigne 1,* , G. Rossbach 1 , J. Levrat 1 , H. Teisseyre 2 , I. Grzegory 2 , R. Butté 1 , T
Suski 2 , and N. Grandjean 1
1 ICMP, Ecole Polytechnique Fédérale de Lausanne, 1015 Lausanne, Switzerland
2 Institute of High Pressure Physics, Polish Academy of Sciences, 01-142 Warsaw, Poland
The recent report of optically pumped III-nitride polariton lasers [1-2] is very promising as it
could lead to a significant decrease of the coherent emission threshold of electrically-driven III-nitride
light-emitting structures thanks to the ultra-low effective mass of polaritons. However, the main
drawback of III-nitride heterostructures, when grown on polar substrates, is the presence of a huge
internal electric field that reduces the exciton oscillator strength and the exciton binding energy in
wide quantum wells (QWs). This can potentially prevent the achievement of the strong light-matter
coupling regime (SCR) leading to the formation of polaritons. In this framework, an alternative
approach consists in developing microcavities (MCs) grown along nonpolar orientations to fully take
advantage of polarization-free QWs.
In the present work, two different MC structures were elaborated by molecular beam epitaxy
using ammonia as nitrogen source on nonpolar m-plane bulk GaN substrates using either a 50 pair
Al 0.15 Ga 0.85 N/Al 0.35 Ga 0.65 N or a 15 pair AlN/Al 0.15 Ga 0.85 N bottom distributed Bragg reflector (DBR).
Five sets of 4 GaN (5 nm)/Al 0.13 Ga 0.87 N (5 nm) QWs placed at the cavity antinodes or a 30 GaN (5
nm)/Al 0.13 Ga 0.87 N (5 nm) multiple QW region are embedded in the 3λ cavity region. Large QW widths
were used to fully take benefit of the absence of built-in electric field: increasing confinement and
enhancing oscillator strength. Finally, a 8 pair SiO 2 /ZrO 2 top dielectric DBR is grown by plasma
enhanced chemical vapor deposition to complete the structures. Room temperature reflectivity has
been first carried out on the bottom DBRs prior to the growth of the active region. As shown in Fig. 1,
the AlN/Al 0.15 Ga 0.85 N DBR is characterized by a flat stopband centered at 3.6 eV with a width of 360
meV. The central DBR wavelength is polarization dependent as a consequence of the strong optical
anisotropy along the nonpolar orientation [3]. The quality of the MQWs was also assessed by
temperature-dependent photoluminescence leading to a narrow linewidth of 60 meV at 300 K (15 meV
at 4 K). In addition, the position of the exciton emission line does match that of the stopband one
ensuring optimum conditions to achieve the SCR. Hints for the strong coupling regime and nonlinear
light emission will be reported.
Fig. 1 : Reflectivity measurements at room temperature performed on
the AlN/Al 0.15 Ga 0.85 N nonpolar DBR as function of polarization angle.
[1] S. Christopoulos et al., Phys. Rev. Lett. 98, 126405 (2007).
[2] G. Christmann et al., Appl. Phys. Lett. 93, 051102 (2008).
[3] T. Zhu et al., Appl. Phys. Lett. 92, 061114 (2008).
Electronic mail: amelie.dussaigne@cea.fr

TuP06
Annealing effects on site-selective InAs quantum dots
M. Helfrich 1,2,* , J. Hendrickson 3 , M. Gehl 4 , D. Rülke 2 , D. Z. Hu 1,2 , M. Hetterich 2 ,
S. Linden 5 , M. Wegener 1,2 , H. Kalt 1,2 , G. Khitrova 4 , H. M. Gibbs 4 and D. M. Schaadt 1,2
1
DFG-Center for Functional Nanostructures (CFN), Karlsruhe Institute of Technology (KIT),
Wolfgang-Gaede-Str. 1a, 76131 Karlsruhe, Germany
2
Institut für Angewandte Physik, Karlsruhe Institute of Technology (KIT), Wolfgang-Gaede-Str. 1,
76131 Karlsruhe, Germany
3
Solid State Scientific Corp., 27-2 Wright Road, Hollis, NH 03049, U. S. A.
4
College of Optical Sciences, University of Arizona, 1630 E. University Bld., Tucson, AZ 85721, U. S. A.
5
Physikalisches Institut, University of Bonn, Nußallee 12, 53115 Bonn, Germany
Nowadays, quantum dot locations can be controlled very precisely through lithographic techniques by
creating small holes on the substrate surface which act as preferential nucleation sites [1,2].
Deterministic assembly of single photon sources or quantum memories based on single quantum dots
is coming closer to realization [3,4]. However, site-selective quantum dots still exhibit inferior optical
properties compared to self-assembled ones [5]. Furthermore, ultimate control of quantum dot
occupation numbers per defined site as well as quantum dot densities has not been achieved. In-situ
annealing is believed to address these problems as it has proven to be a viable tool in manipulating
self-assembled quantum dots [6,7].
Fig. 1: Atomic force microscopy images of site-selective InAs quantum dots as grown (left) and after 2:30 min annealing
(right) [8]. The color scale is xxx nm.
Here we report on our latest results on annealing studies of InAs quantum dots site-selectively grown
on pre-structured GaAs substrates. Small holes were created on the substrate surface by electron beam
lithography prior to molecular beam epitaxial growth of the quantum dots. This is followed by in-situ
annealing with different parameters such as annealing time and temperature being varied. We have
observed a morphological transition of double dots merging into single dots upon annealing, as
illustrated in Fig. 1 and reported in [8]. Atomic force microscopy is used to characterize the
morphology, whereas the optical quality of the site-selective quantum dots is probed in
photoluminescence measurements.
[1] P. Michler et al., Science, 290, 2282 (2000).
[2] N. Akopian et al., Phys. Rev. Lett., 96, 130501 (2006).
[3] S. Jeppesen et al., Appl. Phys. Lett., 68, 2228 (1996).
[4] O. G. Schmidt et al., Surf. Sc. 514, 10 (2002).
[5] T. J. Pfau et al., Appl. Phys. Lett., 95, 243106 (2009)
[6] D. M. Schaadt et al., Appl. Phys. A - Mat. Sci. Proc., 83, 267 (2006).
[7] D. Z. Hu et al., J. Crystal Growth, 312, 447 (2010).
[8] M. Helfrich et al., J. Cryst. Growth, submitted.
__________________________
* Contact: mathieu.helfrich@kit.edu

TuP07
Comparison of InAs quantum dots grown by ripening on InP and
GaInAsP buffer layers on InP(001)
P. Regreny*, A. Benamrouche, C. Brillard and M. Gendry
Université de Lyon, Institut des Nanotechnologies de Lyon, INL-UMR 5270, CNRS, Ecole Centrale de Lyon,
Ecully, F-69134, France
Single self-assembled quantum dots (SAQDs) have attracted a great deal of attention as device
material for quantum information processing such as a single photon emitter for quantum
cryptography [1]. In order to apply QDs to a single photon source for optical fiber-based quantum
cryptography, QDs emitting at 1.3-1.55 µm wavelength band are desirable. InAs QDs on an InP
substrate are promising materials which can emit at the wavelength band due to smaller lattice
mismatch (~3.2 %) permitting larger QDs than other candidates such as InAs/GaAs QDs. Improving
extract efficiency of photon is another challenge for a single photon source. One possible way to
overcome this problem is to confine a single QD in a microcavity such as photonic crystal (PhC)
microcavity. In order to confine only one QD in a cavity, decreasing the density of QDs is required.
Ripening process is an effective way to control InAs QD density on InP buffer layer [2]. In this study,
we compare the ripening process on InP and In 0.8 Ga0 0.2 As 0.43 P 0.57 (Q1.18µm) buffer layers. We grew
InAs QDs by solid source molecular beam epitaxy (SSMBE) in a RIBER Compac21 reactor equipped
with arsenic and phosphorus valved cracker cells. A 200 nm InP or Q1.18µm layer was deposited on
an InP(001) semi-insulating substrate. InAs was grown on the buffer layer at 510 °C with a growth
rate of 0.2 ML/s. The amount of deposited InAs was 0.8 ML. After growth of InAs, ripening was
performed at same temperature under As 2 pressure of respectively 2×10 -6 torr and 3 ×10 -6 torr for
different times. The QDs were formed during this process. We observe very different shapes and
densities as a function of the buffer and time. This behavior will be discussed.
30 QDs/µm
2
13 QDs/ / µm 2 4 QDs/ / µm 2
Fig 1: AFM images of InAs QDs/InP buffer versus ripening time (a) 2min, (b) 4 min, (c) 6 min.
30 QDs/µm 2 110 QDs/µm 2
80 QDs/µm 2
Fig 2 : AFM images of InAsQD/Q1.18 buffer versus ripening time (a) 2min, (b) 4 min, (c) 6 min.
[1] P. Michler, A. Kiraz, C. Becher, W. V. Schoenfeld, P. M. Petroff, L. Zhang, E. Hu, and A. Imamoglu,
Science, 290, 2282 (2000).
[2] E. Dupuy, P. Regreny, Y. Robach, M. Gendry, N. Chauvin, E. Tranvouez, G. Bremond, C. Bru-Chevallier and G.
Patriarche, Appl. Phys. Lett., 89, 1231142 (2006).
__________________________
* Contact: Philippe.regreny@ec-lyon.fr

TuP08
Study on homoepitaxial germanium nanowire growth
J. Schmidtbauer 1,* , R. Bansen 1 , T. Boeck 1 , R. Heimburger 1 , Th. Teubner 1
and T. Schoeder 2
1
Leibniz-Institute for Crystal Growth, Max-Born-Str. 2,12489 Berlin, Germany
2 IHP, Im Industriepark 25, 15236 Frankfurt (Oder), Germany
Germanium has gained renewed interest within recent years for aggressively scaled Si CMOS as well
as Si photonics technologies. The reasons are given by the facts that: (i) Ge is compatible with Si
process technology, (ii) it has higher hole concentration and mobility compared to many III-V
semiconductor materials and (iii) its band gap value matches the wavelength of typical
telecommunication infrastructure. With respect to the later application we investigated germanium
nanowires as promising structures for low-band-gap photonics devices compatible with Si CMOS
process technologies.
We studied the influence of different types of Ge substrates on growth direction, growth rate and shape
of MBE grown Ge nanowires. In particular, homoepitaxial growth of the nanowires was compared
using (100), (110) and (111) oriented germanium substrates . In all cases, the experiments revealed
preferential growth in direction. A clear dependency of the growth rate on the inclination of
nanowires towards the surface normal was shown. This growth behaviour is explained by different
amounts of material that contribute to nanowire growth through direct impingement on the sidewalls
of the nanowire.
In addition, we report on germanium nanowire growth on silicon (111) substrates which were
overgrown with a Ge(111) epilayer. In order to reduce the lattice misfit, a 10 nm thick cub-Pr 2 O 3 (111)
buffer has been deposited by MBE prior to the germanium layer. The resulting Ge on insulator
structures are suitable for epitaxy and thus nanowires could be grown successfully thereon with similar
growth behaviour as on germanium bulk substrates.
____________________________
* Contact: schmidtbauer@ikz-berlin.de

TuP09
Correlating electronic and structural properties of Ga-assisted GaAs
nanowires via cathodoluminescence imaging
J.-S. Hwang, J. Kasprzak*, F. Donatini, C. Bougerol, H. Mariette, Le Si Dang and R. Songmuang
CEA-CNRS Group "Nanophysique et Semiconducteurs,’’Institut Néel/CNRS, BP 166,38042
Grenoble Cedex 9, France
Research on semiconductor nanowires (NWs) is a widely developing field, mainly stimulated
by the potential for creating novel optoelectronic devices. Our interest in the nanowire geometry
is driven by two-fold reasons. Firstly, thanks to their one dimensional nature and sub-wavelength
diameters they can efficiently serve as waveguides and provide a high extraction efficiency of the
guided photons [1]. Secondly, as an alternative to Stranski Krastanow quantum dots (QDs), NWs
offer a novel way to fabricate individual quantum emitters by incorporating short insertions of a
smaller band gap material into a larger band gap NW. While the position and shape of slice-in-awire
QDs can be well controlled within the bottom-up growth [2], their subtle relative energy
tuning can be realized by applying external fields [3]. Above features set NWs as an attractive
host to demonstrate long-range coherent radiative transfer within a pair of individual quantum
emitters [4].
Via this communication, we report our preliminary results on the growth and optical
characterizations of Ga-assisted GaAs NWs [see Fig. 1(a)]. The NWs were grown by standard III-
As molecular beam epitaxy on Si (111) substrates covered with native oxide. The Ga and As
fluxes were simultaneously deposited on the substrate with III/V ratio around 1. Typical
transmission electron microscopy images reveal 4 important crystallographic zones in NWs.
Firstly; the bottom part always consists of zinc blende phase with twin domains. Then, there is a
transition region consisting of a mixture of zinc blende and wurtzite domains with various lengths.
In the third part, a wurtzite region is observed with a few stacking faults. Finally, the top part (~a
few hundred nanometers) closed to the Ga-catalyst shows pure wurtzite phase [Figs.1 (b)-(c)].
We have performed poly- and mono-chromatic low-temperature cathodoluminescence
imaging (CLI) – see Fig.2, giving an insight into the correlation between their structural and
electronic properties. We find that optical (spectral and intensity) response of individual NWs is
highly inhomogeneous. The emission energy is in the range of 1.42-1.48 eV which is not
consistent to the free exciton energy of zinc blende and wurtzite GaAs. This luminescence is
attributed to type II recombination occurring in GaAs “quantum heterostructures”,whose interface
is distinguished by two crystalline phases: zinc blende and wurtzite [5].
The growth control over the wurtzite/zinc blende phase mixing possibly enable fabrication of
slice-in-a-wire individual emitters created by the potential landscape fluctuation of a bicrystalline
heterostructure, in alternative to usual ones based on chemically different
semiconductors.
[1] J. Claudon et al. Nature Photonics 4, 174 (2010).
[2] J. Renard et al., Nanoletters, 8, 2092 (2008).
[3] R. B. Patel et al., Nature Photonics 4, 632 (2010).
[4] G. Parascandolo and V. Savona Phys. Rev. B 71, 045335 (2005).
[5] D. Spirkoska et al., Phys Rev. B 80, 245325 (2009).

TuP09
Fig 1: (a) Scanning transmission electron micrograph (SEM) of Ga-assisted GaAs NWs grown on Si(111) substrate. (b)
Transmission electron microscopy (TEM) image of dispersed NWs indicating Ga droplets on the NW top and a
transition region consist of high density of stacking faults. (c) High resolution TEM of the wurtzite phase close to Gacatalyst.
Fig 2: (a) CL spectrum obtained on an individual GaAs NW.
(b) SEM image of a small ensemble of GaAs NW (left) and corresponding polychromatic CLI (right)
(c) SEM image of an individual GaAs NW (left) and corresponding mono-chromatic CLI at different wavelengths.

TuP10
Growth and microstructure of GaN:Cu
P. R. Ganz 1,2,* , G. Fischer 3 , C. Sürgers 1,3 , H. T. Hsing 4 , L. Chang 4 and
D. M. Schaadt 1,2
1
DFG-Center for Functional Nanostructures, Karlsruhe Institute of Technology (KIT), 76131 Karlsruhe,
Germany
2 Institut für Angewandte Physik, Karlsruhe Institute of Technology (KIT), 76131 Karlsruhe, Germany
3 Physikalisches Institut, Karlsruhe Institute of Technology (KIT), 76131 Karlsruhe, Germany
4 Department of Materials and Optoelectronic Science / Center for Nanoscience and Nanotechnology, National
Sun Yat-Sen University, Kaohsiung 80424, Taiwan, R. O. C.
Diluted magnetic semiconductors (DMS) showing ferromagnetic behavior at room-temperature have
raised interest due to their possible applications in spintronic devices. Especially, DMS based on
semiconducting group-III nitrides are attractive for device applications by reason of their long and
temperature independent spin-lifetime in InN quantum dots [1, 2] and their thermal and chemical
stability. To avoid problems with transition metals precipitates [3] by doping with intrinsic magnetic
elements, Copper was used as dopand. Although Cu is a non-magnetic element, ferromagnetic behavior
was reported for Cu implanted films [4], nanowires [5] and MBE grown films [6]. However, many
questions concerning the low incorporation of Cu into the GaN host and the origin of the ferromagnetic
behavior are not answered yet.
We carried out a detailed study on the growth of GaN:Cu by plasma assisted MBE. The influences of
growth parameters such as metal-to-nitrogen flux ratio on the structural and magnetic properties were
analyzed.
For metal rich growth conditions, the formation of compounds on the surface was observed (Fig. 1a).
X-ray diffraction and transmission electron microscopy indicated the formation of Cu 9 Ga 4 islands. A
cross section through a big island (Fig. 1b) - obtained with a focused ion beam - shows a thick Cu 9 Ga 4
layer epitaxially grown on GaN. The thicker GaN layer underneath the island results from the VLS
(vapor liquid-solid) mechanism. The Cu content in the film was found to be far below the nominal
doping, because of the high amount of Cu localized in the Cu 9 Ga 4 islands.
10 µm
(a)
Fig 1: (a) The surface of GaN:Cu grown under Ga rich conditions. (b) Cross section through an island.
(b)
In order to avoid the formation of Cu 9 Ga 4 islands and to improve the incorporation of Cu into the GaN
host, samples were also grown under nitrogen rich conditions.
[1]S. Nagahara et al., Appl. Phys. Lett. 86, 242103 (2005)
[2]S. Nagahara et al., Appl. Phys. Lett. 88, 083101 (2006)
[3] J. M. Baik et al., Appl. Phys. Lett. 82, 583 (2003)
[4] J.H. Lee et al., Appl. Phys. Lett., 90, 032504 (2007).
[5] H.K. Seong et al., Nano Lett., 7, 3366 (2007).
[6] P.R. Ganz et al., J. Cryst. Growth, in press: doi:10.1016/j.jcrysgro.2010.10.115
__________________________
* Contact: philipp.ganz@kit.edu

TuP11
Different strategies towards the deterministic coupling
of a Single QD to a Photonic Crystal Cavity Mode
J.Herranz 1,* , I.Prieto 1 , Y.González 1 , J.Canet‐Ferrer 2 , P.A.Postigo 1 , B.Alén 1 , L.González 1 ,
L.J.Martínez 1 , M.Kaldirim 1 , D. Fuster 2 , G.Muñoz‐Matutano 2 , and J.Martínez‐Pastor 2
1 IMM-Instituto de Microelectrónica de Madrid (CNM-CSIC), Isaac Newton 8, PTM, E-28760 Tres Cantos,
Madrid, Spain
2 UMDO (Unidad Asociada al CSIC-IMM), Instituto de Ciencia de Materiales, Universidad de Valencia, P.O.
Box 22085, 4607 Valencia, Spain
A single Quantum Dot (QD) coupled to a photonic cavity mode is the fundamental system for
the study of Cavity Quantum Electrodynamics (CQED) phenomena in the solid state
approximation [1], a very promising field for the development of quantum information
technologies. Strong requirements on the simultaneous spectral and spatial matching of the
emitter and the optical cavity mode make the fabrication of single QD-cavity mode coupled
system very challenging. Although coupling of single self-assembled QD to a photonic crystal
(PC) cavity mode has already been demonstrated [2, 3], the technology is far from being
mature. In this sense, the use of high spatial resolution lithographic techniques for site
controlled QD formation[4, 5] is crucial in order to improve the yield of deterministic
integration of a coupled QD – cavity mode[6,7].
In this work we present two strategies for coupling InAs site-controlled QD with the optical
mode of GaAs-based PC microcavities. Site controlled InAs QDs are obtained using AFM
local oxidation lithography to define preferential nucleation sites during MBE growth. These
nanostructures show good optical emission and can be operated as single photon sources [5].
The first approach consists of the fabrication of PC microcavities on GaAs slabs. On top of
the microcavities, AFM local oxidation lithography is performed to define the nucleation site
of a single QD within the cavity. They are transferred into the MBE chamber to complete the
structure to the target thickness with an embedded InAs QD at the predefined site. Our results
show that this process leads to a strong evolution of the round shape of the PC holes, that
degrades the quality factor (Q) of the cavity mode.
In the second approach, same techniques are employed; an etched ruler is fabricated by using
e-beam lithography and dry etching (RIE) on a GaAs epitaxial slab. Then, AFM local
oxidation lithography is performed to define the nucleation sites of InAs QDs. The position
coordinates are set and recorded with respect to the fabricated ruler. The sample is then
completed by a MBE regrowth process. Once it is completed, PC cavity can be designed for
matching the emission wavelength of the embedded QD and later fabricate around the
previously recorded position coordinates of the QD with respect to the ruler. Optical micro-
PL measurements carried out by confocal microscopy at 77K of the obtained nanostructures
in both approaches will be shown.
[1] T. Yoshie et al., Nature 432 (2004) 200.
[2] A. Badolato et al., Science, 308 (2005), 1158.
[3] K. Hennesy et al., Nature 445 (2007), 896.
[4] T. Sünner et al., Opt. Lett 33 (2008), 1759.
[5] J. Martin-Sanchez et al., ACS Nano 3 (2009), 1513.
[6] S. M. Thon et al., App Phys Lett 94 (2009) 111115.
[7] D. Englund et al., Proc. SPIE 7611 (2010) 7611OP.
__________________________
* Contact: jesus@imm.cnm.csic.es

TuP11
Figures
Figure 1. First approximation: Site controlled InAs QD fabrication at the maximum of the electric field of a pre-patterned
photonic crystal cavity. Different process steps and the corresponding AFM images: (a) Fabrication of the photonic crystal
structure (SEM-RIE) and GaAs oxide dot (AFM local oxidation) on a GaAs 105 nm thick slab; (b) The structure is completed
by a MBE regrowth process up to 140nm thickness with an embedded site-controlled InAs QD placed 20nm below the
surface. AFM profiles show the change on the hole shape due to the regrowth step.
a b c
Figure 2. Second approximation: Deterministic cavity fabrication around site controlled InAs QD positioned respect to a prepatterned
etched ruler. (a) Sketch of the final structure with a photonic cavity fabricated at the position of a site-controlled
InAs QD, with coordinates set by the etched ruler. (b) SEM image of the fabricated etched ruler with the alignment marks. (c)
AFM image of an actual GaAs oxide dot obtained by AFM local oxidation with known coordinates respect to the ruler.

TuP12
Capping effect on the morphological and optical properties of
GaAs/AlGaAs quantum structures
M. Jo * , G. Duan, T. Mano and K. Sakoda
National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047,Japan
Droplet epitaxy is a self-assembled growth technique based on the formation of metallic droplets
followed by crystallization into semiconductor quantum dots (QDs). Droplet epitaxy allows the
self-assembly of QDs in lattice-matched systems such as GaAs/AlGaAs, which is unattainable in a
conventional Stranski-Krastanow growth mode. In the growth of GaAs/AlGaAs QDs, various quantum
structures such as monomodal dots, single/multiple rings, and nanoholes have been derived by
controlling the As pressure and temperature during the crystallization of Ga droplets.
However, in droplet epitaxy, low-temperature processes at around 200°C are required for the
formation of droplets and their crystallization, which often causes degradation of the crystalline and
optical qualities of the QDs and subsequent AlGaAs capping layer. Uncapped annealing of QDs is
therefore used as an effective way to improve the quality of the QDs, but this annealing step can also
cause significant morphological changes in the QD.
Here, we report the capping effect on the morphological and optical properties of GaAs/AlGaAs
QDs grown by droplet epitaxy. First, we studied the optical qualities of GaAs nanostructures capped
with a low-temperature AlGaAs layer. To clarify the effects of the capping layer, we used high-quality
GaAs/AlGaAs single quantum wells capped at various temperatures. Luminescence study showed the
formation of abundant nonradiative recombination centers in an AlGaAs capping layer grown at 200°C,
while there is a slight degradation of the optical quality in AlGaAs capping layers grown at temperatures
above 350°C compared to that of a high-temperature capping layer.
Next, we performed thin AlGaAs capping on GaAs QDs to suppress the morphological changes in
the QD. The capping temperature of 400°C was chosen so that no significant morphological change
occurs in the QDs, in addition to the growth of high-quality AlGaAs. AFM measurements revealed that
a thin AlGaAs layer uniformly covers the GaAs QDs. This capping layer provides protections against
thermally induced deformation up to 580°C, which allows improved dot quality. In addition, further
annealing at 640°C flattens the top of the dots, leading to the formation of height-controlled QDs.
.
__________________________
* Contact: JO.Masafumi@nims.go.jp

TuP13
Optical Signatures of Dopant Complexes in Arsenic Doped
HgCdTe Epilayers.
F. Gemain 1* , I. C. Robin 1 , B. Polge 1 and A. Lusson 2
1
CEA-LETI Minatec, 17 rue des martyrs, 38054 Grenoble Cedex 9, France.
2 GEMaC, CNRS, UVSQ, 1 place Aristide Briand, 92010 Meudon, France.
Mercury cadmium telluride (MCT) is extensively used for infrared detector applications. P-type
doping during molecular beam epitaxy (MBE) growth has been widely studied, because it is the key
toward the realization of complex infrared detector heterostructures such as dual band sensors, hot
detectors, and p-on-n-devices [1]. Arsenic is the most used impurity, as it has demonstrated very low
diffusion properties into HgCdTe and also because the available purity and associated effusion
technology meet today’s requirements for MBE. Arsenic can thus be incorporated in levels ranging
from 1.10 16 to 1.10 19 at / cm 3 but the acceptor behavior can only be achieved after high temperature
thermal activation. The incorporation of dopants as well as the different annealing procedures used for
the activation of the dopants lead to significant changes of the luminescence properties [2]. Thus,
photoluminescence (PL) measurements are found to be an efficient tool to control arsenic doping in
the material and its evolution under different post-growth annealings.
For this study, 2 MBE grown samples with a cadmium composition of 47% and 30% were
investigated. Excitation-power dependent and temperature dependent comparative studies were done
to identify the nature of the PL peaks. A piece of each sample underwent a p-type annealing which
consists in a thermal activation at 370°C for 1h under saturated mercury pressure to transfer the
incorporated As onto the proper sites [3]. Another piece was n-type annealed which consists in a five
days annealing at 270°C in order to fill the mercury vacancies responsible for uncontrolled p-type
doping. A third piece underwent a p-type annealing plus a n-type annealing to fill the mercury
vacancies generated during the p-type annealing.
Modulated PL measurements with a continuous-scan Fourier transform infrared spectrometer were
performed between 2K and 300K using a 1064 nm wavelength of a Nd YAG laser for excitation. The
signal was detected with a cooled InSb detector.
The PL study of As doped HgCdTe samples with a cadmium composition of about 30% was correlated
to extended x-ray absorption fine structure (EXAFS) results [6]. The EXAFS results showed the
sharing of incorporated As atoms during MBE growth between two structures: a donor complex
As 2 Te 3 and an AsHg acceptor complex. After p-type annealing, the As 2 Te 3 is dissociated; the amount of
AsHg acceptors is re-enforced and a donor As Hg (As atom into an Hg site) is activated.
Those EXAFS results allowed us to identify the PL peaks in the case of the samples with a cadmium
concentration of 30%. Those results will be used to identify as well the optical signatures of the
different complexes in the sample with a cadmium concentration of 47%.
[1] J. Baylet, P. Ballet, P. Castelein, F. Rothan, O. Gravrand, M. Fendler, E. Lafosse, J. P. Zanatta, J. P. Chamonal, A.
Million, and G. Destefanis, J. Electron. Mater. 35, 1153 (2006).
[2] I. C. Robin, M. Taupin, R. Derone, A. Solignac, P. Ballet, and A. Lusson, Appl. Phys. Lett., 95, No. 202104 (2009).
[3] P. Ballet, B. Polge, X. Bicquard, and I. Alliot, J. Electron. Mater. 38, 1726 (2009).
__________________________
* Contact: frederique.gemain@cea.fr

TuP13
As2Te3
AsHg
Band to band
AsHg
Figure 1 Comparison of the PL spectra at 2K of the as grown sample, after n-type n
annealing and after p-type p
plus n-type n
annealing for the sample with a Cd
concentration of 30%. The peaks were identified using the EXAFS results.

TuP14
Optical properties of post-growth annealed type-II GaSb
quantum dots
A. Schramm * , V. Polojärvi, T. V. Hakkarainen, A. Gubanov, J. Paajaste, R.
Koskinen, S. Suomalainen, and M. Guina
Optoelectronics Research Centre, Tampere university of Technolgoy,
P.O. Box 692, FIN-33101 Tampere, Finland.
Semiconductor quantum-dot (QDs) ensembles obtained from Stranski-Krastanov growth have
attracted great interest during the past few decades because of their applications in novel devices such
as QD lasers, solar cells, or photodetectors as well as memory devices. While the commonly exploited
InAs/GaAs QD systems exhibit a type-I band alignment, GaSb/GaAs has a type-II band structure, i.e.
the holes are confined while the non-confined electrons are loosely bound by Coulomb interactions
leading to spatial charge separations. The main advantages of using GaSb QDs is related to the
availability of a wide choice of barriers materials that can be used for band-gap engineering of
telecomm emitters [1] or solar cells [2].
Here, we study optical properties of post-growth annealed GaSb QDs grown by molecular beam
epitaxy (MBE) with different V/III ratios in the Stranski-Krastanov regime. Increasing the V/III ratio
yields larger QDs, decreases the photoluminescence (PL) intensity, and shifts the photoluminescence
to lower PL energies. Post-growth annealing below and above ~800 ºC increases the PL intensities and
causes PL blueshifts, respectively. The relative blueshift is similar in both samples indicating
intermixing processes independent on V/III ratio and thus, on composition and size. The optical
properties of the as-grown and annealed QD samples are studied by means of temperature, excitation
power density and polarization dependent PL measurements.
PL signal (arb. u.)
1.0
0.8
0.6
0.4
0.2
0.0
(a)
QDs
V/III = 3
1.0 1.1 1.2 1.3 1.4 1.5 1.6
E (eV)
V/III = 7
WL GaAs
PL signal (arb. u.)
1.6
1.2
0.8
0.4
0.0
as-grown
700 o C
800 o C
V/III = 7
900 o C
950 o C
1000 o C
1.0 1.2 1.4
E (eV)
(b)
V/III = 3
1.0 1.2 1.4
E (eV)
(c)
Fig. 1. (a) Room-temperature (RT) PL spectra of GaSb QDs grown with V/III ratios of 3 and 7. The insets show atomic-force
microscopy pictures of corresponding surface QD samples. In (b) and (c) are RT-PL spectra of annealed samples with V/III
ratios 7 and 3, respectively.
[1] S. Lin et al. , Appl. Phys. Lett. 96, 123503 (2010),
[2] R. B. Laghumavarapu et al., Appl. Phys. Lett. 90, 173125 (2007).
__________________________
* Contact: andreas.schramm@tut.fi

TuP15
Metamorphic 6.3Å GaInSb templates grown on GaAs substrates
for mid-infrared lasers
L.Cerutti * , J.B. Rodriguez and E. Tournié
Institut d'Electronique du Sud, Université Montpellier 2, UMR CNRS 5214,
Place Eugène Bataillon, CC 067, F-34095 cedex 5 Montpellier (France)
The 2-4 µm wavelength range presents a great interest for spectroscopy due to the presence of
absorption lines of several pollutants such as CO 2 , HF, CH 4 ,... Nowadays, the 2-3 µm wavelength
range is covered with lasers grown on GaSb substrates and made of type I GaInAsSb quantum wells
(QWs) embedded in AlGaAsSb. To reach longer wavelength, the AlGaAsSb quaternary barrier has to
be replaced by the AlGaInAsSb quinary alloy in order to keep the type I alignment. However, the
reproducibility of the crystalline perfection and the lattice matching of this quinary alloy on GaSb is
difficult to control. In order to avoid the use of this complicated material, we propose a new approach
based on the (Al,Ga,In)Sb material. Indeed, the AlGaInSb material allows realizing type I QWs for
emission in the 3-4 µm range with a single group-V element. The major problem is that the
corresponding alloys cannot be lattice-matched to any substrate. However, it has been shown that
under adequate conditions the Sb-based materials grown on highly lattice-mismatched substrates relax
the strain mainly by formation of a self arranged 2D array of dislocations confined at the heterointerface
[1,2]. This results in a reduction of the threading dislocations density which improves the
crystalline perfection of the materials without the need to grow a thick buffer layer. In this context, we
have chosen to use AlGaInSb with a lattice parameter around 6.3Å. This allows covering a large range
of energy gap (from 0.32eV to 1.2 eV) and optical index and opens the way to realize high
performance MIR lasers. Moreover, the growth will be realized on GaAs substrates to gather
numerous advantages (low cost, high quality and large wafer scale).
In order to achieve high performance metamorphic devices, it is important to first grow a high
quality template. In this work we have studied the MBE growth and properties of Ga 0.42 InSb buffer
layers. The samples were grown in a RIBER C21E equipped with valved cracker-cells for group-V
elements. The growth temperature was measured via the calibrated substrate-oven thermocouple. After
de-oxidation of the GaAs substrate under As flux, a 200-nm GaAs buffer layer was grown at 580 °C.
Then, the As- valve and shutter were closed and the sample was soaked in an Sb flux before
decreasing the substrate temperature. Under this condition, the RHEED pattern shows a 2x8
reconstruction related to a Sb-stable GaAs surface [3]. The 500-nm GaInSb template was then grown
at 1ML/sec with a V/III growth rate ratio of 2. All layers have been characterized by AFM, HRXRD
and PL.
Figure 1 shows the evolution of the HRXRD pattern with the growth temperature T g . When T g
exceeds 450°C, the GaInSb peak shifts toward the GaAs substrate peak indicating a reduction of the In
content. Indeed, at temperatures higher than 450°C, we approach the melting point of InSb (525°C)
which probably reduces the sticking coefficient of In. Moreover, at 500°C the peak is very broad and
we can clearly see 3 peaks revealing a phase separation in the GaInSb alloy. The surface of such
sample is then very degraded. From these characterizations we deduce an optimum growth
temperature of 425°C which is confirmed by the AFM study of the surface RMS roughness (Fig. 2).
Room-temperature PL measurements also show that the highest peak intensity is obtained for buffer
layers grown at 425°C (Fig. 3).
We are currently investigating the epitaxy of active regions on templates grown at 425°C for
emission in the 3-4µm range. The results will be reported at the workshop.
[1] A. Rocher and E. Snoeck, Mat. Sc. and Eng. B67 (1999).
[2] J.B Rodriguez, L.Cerutti, P. Grech and E. Tournié, Appl. Phys.Lett., 94, 023505 (2009).
[3] J.E. Bickel, N.A. Modine, J. Mirecki Millunchick, Surface Science, 603 (2009)
__________________________
* Contact: cerutti@univ-montp2.fr

TuP15
Intensity (a.u.)
10 17
10 15
GaInSb
10 13
10 11
10 9
10 7
10 5
10 3
10 1
375°C (410")
400°C (392")
425°C (370'')
450°C (372'')
475°C (384'')
500°C
10 -1
28 29 30 31 32 33
(°)
GaAs
Fig 1: HRXRD of GaInSb buffer grown on GaAs at different growth temperature
2.0
1.8
2µm x 2µm
5µm x 5µm
20µm x 20µm
RMS (nm)
1.6
1.4
1.2
1.0
0.8
0.6
0.4
370 380 390 400 410 420 430 440 450 460 470 480 490 500
Growth temperature (°C)
Fig 2: RMS roughness for different surface of GaInSb buffer grown at different temperature
PL intensity (u.a.)
11
10
9
8
7
6
5
4
3
2
1
PL 4.5V (DC50%)
cal 2mV
RT
T PSCT
375°C
400°C
425°C
450°C
475°C
0
2.00 2.25 2.50 2.75 3.00 3.25 3.50 3.75 4.00 4.25 4.50 4.75 5.00
Wavelength (µm)
Fig 3: RT PL of the GaInSb buffer grown on GaAs at different growth temperature

TuP16
Reflection high-energy electron diffraction φ scans for the in-situ
monitoring of the growth of GaN nanowires on Si
P. Dogan * , O. Brandt, L. Geelhaar, and H. Riechert
Paul-Drude-Institut für Festkörperelektronik, Hausvogteiplatz 5–7, D-10117 Berlin,Germany
The planar growth of GaN on foreign substrates such as Si leads to a relatively high defect density and
strain due to the lattice and thermal mismatch. An alternative approach of producing GaN films on
foreign substrates is based on the pendeoepitaxial overgrowth and coalescence of GaN nanowires
(NWs) [1]. In order to minimize the threading dislocation formation during the coalescence process,
the in-plane (twist) and out-of-plane (tilt) orientation distribution of the NWs need to be minimized.
We have investigated in-situ both twist and tilt of GaN NWs grown on different substrates, namely
Si(100), Si(111) and AlN-buffered Si(111) by plasma-assisted molecular beam epitaxy. This in-situ
analysis is based on reflection high-energy electron diffraction (RHEED) φ scans and stationary
RHEED images. RHEED φ scans have been used for the determination of surface reconstructions by
Braun et al. [2], as well as for revealing the orientation-relationship of heteroepitaxial structures by
Gao et al. [3]. To perform the scan, RHEED intensities are recorded along a line parallel to the shadow
edge during substrate rotation. This powerful method enables the real-time monitoring of the epitaxial
growth. In particular, the twist of GaN NWs can be monitored by the RHEED φ scans, whereas the tilt
of the NWs is observed in the stationary RHEED images as an arc. We show that in-situ RHEED φ
scans and the stationary RHEED images which are acquired within a couple of minutes offer the same
structural information which otherwise requires time-consuming ex-situ x-ray diffraction (XRD)
measurements. Moreover, these techniques enable the monitoring of the evolution of twist and tilt
during the growth.
Figure 1 (a), (b), and (c) displays the RHEED φ scans of the GaN NWs grown on Si(100), Si(111), and
AlN-buffered Si(111), respectively. The hexagonal symmetry of the GaN(0001) surface can be clearly
distinguished in Fig. 1 (b) and (c) revealing the epitaxial growth of the NWs on these substrates. On
the other hand, the broad distribution observed in Fig. 1 (a) shows that GaN NWs grown on Si(100)
loose their epitaxial relation to the substrate. The width of the reflections along φ gives information
about the twist of the GaN NWs. The analyses reveal that GaN NWs grown on Si(100) show the
largest twist which is beyond the resolution limit of this technique. On the other hand, the twist of the
NWs decreases dramatically for the GaN NWs grown on Si(111) to 4.7° and on AlN-buffered Si(111)
to 3.8°. For comparison, Fig. 2 (a) displays the XRD ω scans recorded for the GaN(201) reflection of
the same samples. As can be seen from the width of the peaks, XRD ω scans agree well with the
RHEED φ scans exhibiting the same trend for the twist. The tilt of the GaN NWs has been monitored
by the stationary RHEED images [Fig. 1 (d), (e), and (f)] and compared with the XRD ω scans
recorded for the GaN(002) reflection [Fig. 2 (b)]. The tilt of the NWs as observed as an arc in the
images reduces dramatically for GaN NWs grown on AlN-buffered Si(111) in comparison to the NWs
grown on Si(111) and Si(100). The same trend of tilt is shown in Fig. 2 (b). GaN NWs grown on AlNbuffered
Si(111) reveal a tilt below 1°, whereas the tilt increases to about 3° for NWs on Si(111) and to
about 5° for NWs on Si(100). In addition, the evolution of the twist and tilt is monitored during the
growth of the NWs and it is found that both twist and tilt of the NWs decrease with the growth time.
These results show that RHEED φ scans and RHEED stationary images can be used as a control tool
during the growth to collect information about the epitaxial quality of the NWs.
[1] K. Kusakabe et al., Jpn. J. Appl. Phys., 40, L192 (2001).
[2] W. Braun et al., J. Vac. Sci. Technol. B., 16, 1507 (1998).
[3] C.X. Gao et al., Appl. Phys. Lett., 97, 031906 (2010).
________________________
* Contact: pinar.dogan@pdi-berlin.de

TuP16
(a) (b) (c)
(d) (e) (f)
Fig 1: RHEED φ scans along a line across the specular spot and RHEED stationary images along the [ 1120]
azimuth for
GaN NWs grown on (a), (d) Si(100), (b), (e) Si(111), and (c), (f) AlN-buffered Si(111), respectively. Note that the hexagonal
symmetry is slightly distorted in (b) due to the shift of the specular spot during recording.
Fig 2: XRD ω-scans of GaN NWs on Si(100), Si(111), and AlN-buffered Si(111) for the (a) GaN(201) (twist) and (b)
GaN(002) (tilt) reflections.

TuP17
Optical polarization from self-organized InP QDs grown on an
self-undulated template
A. Ugur 1 , F. Hatami 1,* N. Vamivakas 2, L.Lombez 2, M. Atatüre 2 and W. T.
Masselink 1
1
Institut f ür Physik der Humboldt-Universität zu Berlin, Newtonstr. 15, D–12489 Berlin,
Germany
2
Cavendish Laboratory, University of Cambridge, JJ Thomson Avenue, Cambridge CB3
0HE, UK
Controlling the position of quantum dots (QDs) is particularly important for device applications, for
instance in lasers [1] or in computing [2]. Sophisticated techniques have recently been developed to
obtain self-assembled highly ordered QD structures through either a stacking layer of QDs and/or
surface miscut [3]. Recently we have demonstrated a method wherein ordered chains of InP/
In 0.48 Ga 0.52 P QDs are formed during their self-organized growth on an undulating In 0.48 Ga 0.52 P buffer
layer [4]. Although it was apparent that the undulations in the [-1 1 0] direction act as a template for
dot growth, two mechanisms — lateral composition modulation (LCM) and long range ordering
(LRO) — were suggested as mechanisms.
We have investigated the polarization of the optical emission from these ordered QDs.
Photoluminescence measurements show enhanced linear polarization in the alignment direction of
QDs, [-110]. These measurements indicate that the mechanism for the undulations and, thus, also for
the ordering of the QDs is LCM.
Micro-photoluminescence from this system shows emission from both the InGaP buffer at 630 nm and
from the InP QDs at 660 nm; both emission bands are optically polarized in the [-1 1 0] direction. Our
measurements show a 46% degree of polarization in undulated InGaP buffer layers, compared to only
16% in non-undulated comparison samples. The degree of polarization from the ordered QDs grown
on the undulated buffers is 66% compared to 36% from non-ordered QDs. All degrees of polarization
are independent of the polarization of the excitation light. The greater degree of polarization from the
undulated InGaP buffer may be explained by lateral compositional modulation (LCM) where the
degeneracy at the top of valance band is lifted, favoring optical transitions polarized along [-110]
direction [5]. The enhanced QD polarization from the ordered QDs on the undulated buffers could be
due either to the additional strain resulting from the undulations or to the valence band splittings
associated with LCM. Radiative recombination in these states may exhibit differing polarizations
according to the appropriate selection rules. Fine splitting of the degenerate quantum dot states into
components containing [110]-allowed, [-110]-allowed, and [001]-allowed optical transitions, will be
discussed.
[1] K. Meneou, K.Y. Cheng, Z.H. Zhang, C.L. Tsai, C.F. Xu, and K.C. Hsieh, Appl. Phys. Lett. 86, 153114 (2005)
[2] T. v. Lippen, R. Nötzel, G.J. Hamhuis, and J.H. Wolter, Appl. Phys. Lett. 85, 118 (2004).
[3] D. E. Wohlert, K. Y. Cheng, K.L. Chang, and K.C. Hsieh, J. Vac. Sci. Technol. B17, 1120 (1999).
[4] A. Ugur, F. Hatami, M. Schmidbauer, Michae Hanke, W.T. Masselink, J. Appl. Phys. 105, 124308
(2009).
[5] U. Hakanson, V. Zwiller, M. K.-J. Johansson, T. Sass, and L. Samuelson, Appl. Phys. Lett. , 82, 627 (2003)

TuP17
Fig.1 (a) AFM micrograph of a sample without cap layer with excitation directions, the height contrast is between 0–6 nm, the scale for
AFM micrograph is 3 µm × µm, the scale for inset micrograph is 0.5 µm × µm, (b) PL spectrum for laterally ordered QDs, emission at 630
nm stems from InGaP and emission at 660 nm stems from InP QDs. (figures are submitted to APL)
FIG. 2. Polarization dependent PL measurements for the samples with QDs, (a) The collection polarization dependence of the InGaP buffer
layer, in sample A, for θe = 0°, 45°and 90°, (b) The collection polarization dependence of the InP QDs, in sample A, for θe = 0°, 45°and 90°,
(c) The collection polarization dependence of the InGaP buffer layer, in sample B, for θe = 0° and 45°, θe = 0° is normalized at 10◦ collection
polarization since the data points for θe = 0° and 5° were noisy, (figures are submitted to APL).
* Contact: ugur@physik.hu-berlin.de
_________________________

TuP18
The MBE growth of HgCdTe on CdZnTe and CdTe/Ge at CEA-
LETI
Giacomo Badano, Philippe Ballet, Sebastien Renet, Philippe Duvaut, Xavier Baudry,
Bernard Polge and Alain Million 1
1
CEA-LETI Minatec, 17 rue des Martyrs, 38054 Grenoble Cedex 9, France
Molecular Beam Epitaxy (MBE) has long been recognized as the most flexible
growth technique for Hg1-xCdxTe. It is the only technique that allows for in-situ doping
and the growth of complex multilayer structures for the next generation of infrared
detectors. Hg1-xCdxTe is second only to Si and GaAs for technological importance. This
alloy dominates most of the high-end infrared detection market thanks to the tunability
of its bandgap, obtained by varying the Cd concentration. After pioneering [1] the growth
of HgCdTe by MBE in the 1980’s, the CEA has pursued continuous improvement of this
technique.
After reviewing the earliest achievements, we will focus on several representative
topics currently researched at our laboratories. They include the reduction of
dislocations and surface defects, growth on large area CdTe/Ge substrates, and the
control of in-situ p-type doping.
Surface defects and dislocations hinder device performance and must be reduced.
In spite of much effort, defects still pose a major challenge. The current state-of-the-art
is on the order of 1000 defects/cm 2 . To reduce it further, better comprehension and
control of growth and strain effects will be needed. The potential of in- and ex-situ
characterization techniques in this respect will be discussed. We will also talk about the
principal types of defects, their origin and impact on device performance.
In a drive to increase availability and reduce cost, large area CdTe/Ge substrates
were introduced in the 1990’s as an alternative to traditional bulk CdZnTe substrates,
which are small and very expensive. The major challenge of CdTe/Ge is its dislocation
density, which has recently been reduced to 5x10 6 cm 2 . Details of their growth,
morphology and relaxation will be given.
Another hot topic in the field is in-situ p-type doping based on arsenic. After
incorporating it into the material, arsenic is thermally activated through a high
temperature annealing. Using EXAFS, the underlying mechanism for arsenic activation
is found to be more complex than a simple change in lattice site. It rather involves a noncrystalline
phase in the vicinity of arsenic [2,3].
Finally, the discussion of the characteristics of recent FPAs obtained from our
layers will illustrate the maturity of our MBE technology.
[1] 1- J.P. Faurie and A. Million, J. Cryst. Growth 54, 582 (1981)
[2]- P. Ballet, B. Polge, X. Biquard and I. Alliot, J. Electron. Mater. 38, 1726 (2009)
[3]- X. Biquard, I. Alliot and P. Ballet, J. Appl. Phys. 106, 103501 (2009)
__________________________
Contact author: giacomo.badano@cea.fr

nitrogen flow [sccm]
growth rate [nm/min]
growth rate [nm/min]
TuP19
Control of nitrogen plasma for growth of GaN by
plasma-assisted MBE
M. Sobanska 1 , K. Klosek 1 , Z.R. Zytkiewicz 1,* , H. Teisseyre 1 , A. Wierzbicka 1 , E.
Lusakowska 1 , and W. Jung 2
1 Institute of Physics, Polish Academy of Sciences, Al. Lotnikow 32/46, 02 668 Warszawa, Poland
2 Institute of Electron Technology, Al. Lotnikow 32/46, 02 668 Warszawa, Poland
GaN layers were grown at temperature of 720 o C on a GaN/sapphire templates using a plasma assisted
MBE in Riber Compact 21 system, equipped with elemental sources of Al, Ga, In, Si, and Mg. Active
nitrogen was supplied from an Addon RF plasma source. An optical Si sensor was used to measure
the light intensity emitted by the nitrogen source in the wavelength range of 750 to 850 nm. The
sensor converts light emitted by excited nitrogen species into output voltage U.
Nitrogen plasma intensity, and so the optical sensor output voltage U, depends both on the RF power
and the nitrogen gas flow. This is illustrated in Fig. 1 that shows quite broad range of RF power and
nitrogen flow parameters giving fixed value of sensor output U. In Fig. 2 the GaN growth rate vs.
nitrogen flow is plotted for U=1.5 V and U=2.4 V. For each point the gallium flux was adjusted to
keep the surface covered by 2ML of Ga. As seen, within our experimental error the growth rate does
not change when the plasma parameters vary along the red (U=1.5 V) and the orange (U=2.4 V)
curves in Fig. 1. On the contrary, the growth rate monotonically increases with intensity of light
emitted by nitrogen plasma as shown in Fig. 3. Since under gallium-rich conditions growth is
controlled by nitrogen flux, results shown in Figs. 2-3 clearly indicate that the optical sensor output
signal is a direct measure of the amount of active nitrogen species available for growth. This
conclusion is similar to that reported earlier by Foxon et al. [1] for plasma-assisted MBE growth of
AlAsN and GaAsN alloys.
7
6
5
4
3
2
1
A
B
C
300 350 400 450 500 550 600
RF power [W]
U = 1.5 V
U = 1.8 V
U = 2.0 V
U = 2.2 V
U = 2.4 V
U = 2.6 V
Fig. 1: Set of RF power and nitrogen
flow parameters for a fixed value of
optical sensor output U.
6,0
5,5
5,0
3,5
3,0
2,5
2,0
U = 2.4 V
U = 1.5 V
1,0 1,5 2,0 2,5 3,0 3,5 4,0
nitrogen flow [sccm]
Fig. 2: GaN growth rate vs. nitrogen
flow for U=1.5 V and U=2.4 V.
6,5
6,0
5,5
5,0
4,5
4,0
3,5
3,0
2,5
N 2
flow = 2.5 sccm
1,4 1,5 1,6 1,7 1,8 1,9 2,0 2,1 2,2 2,3 2,4 2,5 2,6 2,7
sensor output signal U [V]
Fig. 3: GaN growth rate vs. sensor
output voltage U for nitrogen flow
of 2.5 sccm.
Three 0.7 µm thick GaN layers were grown with nitrogen plasma parameters as marked by A, B, and
C in Fig. 1 in order to study if their properties, as seen by atomic force microscopy, X-ray diffraction,
low temperature photoluminescence, and by electrical measurements, depend on the RF power and/or
the nitrogen flow used during growth.
This work was partially supported by the European Union within European Regional Development
Fund, through grant Innovative Economy (POIG.01.03.01-00-159/08 and POIG.01.01.02-00-008/08).
[1] C.T. Foxon et al., J. Vac. Sci. Technol. B, 14, 2346, (1996).
__________________________
* Contact: zytkie@ifpan.edu.pl

TuP20
GaN/AlN semipolar quantum dots for ultra-violet emission
A. Kahouli 1,2 *, N. Kriouche 1 , J. Brault 1 , B. Damilano 1 , P. de Mierry 1 , A. Courville 1 , J.
Massies 1
1. Centre de Recherche sur l’hétéro-Epitaxie et ses Applications, Centre National de la Recherche Scientifique
Valbonne 06560
2. Université de Nice Sophia Antipolis, Parc Valrose, 28 avenue Valrose, 06108 Nice cedex 2, France
GaN quantum dots (QDs) can be of interest to improve the performances of short wavelength
optoelectronic devices. One interesting property is that carriers are confined in the three spatial
dimensions, which leads to the minimization of non-radiative carrier recombinations on crystalline
defects. When GaN/AlN QDs are grown along the [0001] axis of the wurtzite structure (“polar” axis),
a strong internal electric field, induced by the polarization difference at the GaN-AlN interfaces, is
created in the heterostructure. This field induces a strong reduction in QD transition energies, which
offers the possibility to cover the whole visible spectrum [1]. On the other hand, it is more difficult to
obtain an ultra-violet emission. To reduce this quantum confined stark effect, it is possible to grow the
QDs on the so-called “semi-polar” (11-22) orientation [2] (in this case the [0001] axis makes an angle
of 32° with the growth surface).
The GaN/AlN QD structures are grown by molecular beam epitaxy using ammonia on GaN
(11-22) templates obtained by metal-organic vapour phase epitaxial growth on (1-100) sapphire [3].
On these templates, 360 nm of AlN followed by three QD planes of nominally 6 GaN monolayers
buried by 30 nm of AlN are grown. An additional QD plane is deposited on the surface in order to
study the QD morphology by atomic force microscopy (figure 1). A sample with a comparable
structure was grown along the [0001] direction in order to be used as a reference.
We have studied the photoluminescence properties of the semipolar QDs as a function of
temperature and of the excitation power density. The most important result is that, due to the reduced
electric field, the semipolar QD emission is shifted towards higher energies by almost 1 eV compared
to polar QDs (figure 2). An ultraviolet emission at 333 nm is then achieved with those “semipolar”
QDs.
Polar QDs
PL intensity (arb. u.)
x10
GaN template
Semi-polar QDs
300 350 400 450 500 550 600
Wavelength (nm)
Figure 1: Atomic Force Microscopy image of semipolar
GaN/AlN QDs. The image scale is 1x1 µm 2 (derivative
mode)
Figure 2: 14 K PL spectra of polar and semipolar GaN/AlN QDs.
The positions of the QD emission peaks are indicated by vertical
dashed lines.
[1] B.Damilano et al., Applied Physics Letters 75, 962 (1999)
[2] L.Lahourcade et al, Applied Physics Letters 94, 111901 (2009)
[3] P. de Mierry et al, Japanese Journal of Applied Physics 48, 031002 (2009).
________________
* Contact: ak@crhea.cnrs.fr

TuP21
MBE growth approaches for improving Sb-based In 0.5 Ga 0.5 As(Sb)/GaAs QDs.
M.J. Milla 1 , Á. Guzmán, J.M. Ulloa, A. Hierro
Instituto de Sistemas Optoelectrónicos y Microtecnología (ISOM), Dept. de Ingenieria Electrónica, ETSI de
Telecomunicación, Universidad Politécnica de Madrid, Ciudad Universitaria s/n, 28040 Madrid (Spain)
During the last years, the development and improvement of InAs quantum dots (QD) lasers emitting at
1.3 - 1.55 m with a higher efficiency and lower cost is being a subject of intense research. In fact,
different alternatives such as GaInAsN QDs [1] or strain reduction layer (SRL) based on GaAsSb or
InGaAs [2] are being considered. Another possibility to reduce the strain is to employ InGaAs QD in
the active region, adding other elements to shift the emission to longer wavelengths. In particular, the
role of Sb has been widely studied in the case of diluted nitrides quantum wells, where it is observed
that it helps to redshift the wavelength emission, to reduce the number of point defects and the
composition modulation [3]. However, there are only a few publications reporting the effect of Sb in
the optical and morphological properties of QD incorporating Sb into InGaAs QD. In this work, we
study the influence of the presence of Sb during the formation of the QD and the steps after this
formation using different MBE growth procedures. We correlate the optical properties of buried QD
capped with GaAs, with the morphological aspects of similar layers grown on the surface of the
samples.
The addition of antimony during the MBE growth was carried out following different approaches.
One consisted on keeping the Sb open only during the formation of the QD, a second alternative
employed, was to maintain a flux of antimony 10 seconds after the QD formation. Also, we used this
latter procedure but closing the arsenic shutter 20 second after the QD formation, keeping the Sb open.
As a consequence, we observed that the presence of arsenic after the QD formation seems to play an
important role in the final quality of the samples. We attribute this behaviour to the competition
between the two group V elements present during the steps after the QD formation. The different
binding energies among the elements composing the QD are analyzed and compared to establish their
possible effect in the final composition of the QD after the capping process.
We perform an analysis of the influence of the mentioned growth approaches and the Sb flux in the
optical properties of the QD. We obtain an enhancement of PL intensity for samples with a higher
amount of antimony (samples labeled as a in fig.1). Besides, we find that the growth approach plays a
role in the shift of the wavelength (see figure 1): when the arsenic is open after the quantum dot
formation, there is a blueshift [4] of the main peak compared with a pure InGaAs QD reference.
However, if the arsenic is closed, the emission wavelength is not shifted. This could be related with
the possible formation of a Sb coverage on top of the QD which is affected by the presence or not of
the As species [5]. We also observed that samples grown closing the As, present a higher homogeneity
(fig.2).
[1] M. Sopanen, H. P. Xin, and C. W. Tu, Applied Physics Letters, 76, 994 (2000).
[2] J. M. Ulloa, I. W. D. Drouzas, P. M. Koenraad, D. J. Mowbray, M. J. Steer, H. Y. Liu, and M. Hopkinson. Applied
Physics Letters, 90, 213105 (2007).
[3] F. Ishikawa, Á. Guzmán, O. Brandt, A. Trampert, and K. H. Ploog. Journal of Applied Physics, 104 (11), 113502 (2008).
[4] T. Matsuura, T. Miyamoto, T. Kageyama, M. Ohta, Y. Matsui, T. Furuhata and F. Koyama, Japanese journal of applied
physics, 43(5A), L605 (2004).
[5] S.I. Molina, A.M. Sánchez, A.M. Beltrán, D.L. Sales,T. Ben, M.F. Chisholm, M. Varela, S. J. Pennycook, P.L. Galindo,
A.J. Papworth, P.J. Goodhew, J.M. Ripalda, Applied Physics Letters, 91, pp. 263105 (2007).
1 Corresponding author: María José Milla (mjmilla@die.upm.es)

TuP21
PL intensity (a.u.)
reference = InGaAs QD
a Sb BEP = 1.1·10 -7 mbar
b Sb BEP = 4.7·10 -8 mbar
b
a
b
a
reference
As openned
As closed
950 1000 1050 1100
Wavelength (nm)
Figure 1. Influence in the optical properties of different growth approaches and Sb fluxes.
Figure 2. ( 1µmx1µm) AFM images of InGaAsSb QDs . Vertical scale is 7nm and surface density 10 11 cm -2

TuP22
Post-growth rapid thermal annealing of InAs quantum dots
grown on GaAs nanoholes formed by droplet epitaxy
L. Wewior, B. Alén, D. Fuster, L. Ginés, Y. González, J. M. Llorens, D. Alonso-
Álvarez, and L. González
IMM-Instituto de Microelectrónica de Madrid (CNM-CSIC), Isaac Newton 8, 28760 Tres Cantos, Spain
The deposition of InAs on GaAs nanoholes formed by droplet epitaxy can be used to tailor the
shape and size of InAs QDs while preserving the low areal density (~2x10 8 cm -2 ) imposed by the
nanohole pattern [1,2]. Our previous micro photoluminescence (µPL) study of individual InAs QDs
grown by this method revealed that the single QD emission was largely affected by the charged
environment surrounding the nanostructure. This charged environment, attributed to the presence of As
vacancies, leads to multicharged exciton emission and spectral diffusion effects which might limit the
suitability of these QDs in quantum light emitting applications. An intense single peaked emission
with radiation limited linewidth and null fine structure splitting (FSS) would be desirable to fully
exploit the size and shape control capabilities of droplet epitaxy based methods. The aim of the present
study is to reduce the presence of the As vacancies near the QDs by applying a post-growth rapid
thermal annealing (RTA) to the sample. Besides, the small FSS (~41 µeV) [1] found typically in these
QDs might be reduced further by the RTA processes [3].
PL I nt e ns i t y ( a r b. uni t s )
P L I n t e n s i t y ( a r b . u n i t s )
P e a k : 1 . 2 9 7 e V P e a k : 1 . 3 0 7 e V
F W H M : 1 4 . 7 m e V F W H M : 1 4 . 6 m e V
τ = 6 3 3
τ = 5 2 6
1 . 2 2 1 . 2 4 1 . 2 6 1 . 2 8 1 . 3 0 1 . 3 2 1 . 3 4
1.28 1.29 1.30 1.31
Fig. 1. Time integrated PL spectrum and time resolved PL
curves (inset) before and after the RTA process.
p s
A s g r o w n
5 mi n RT A
( 7 7 5 º C)
p s
4 6 8 1 0 1 2 1 4 1 6 1 8 2 0
T i m e ( n s )
p-shell
P h o t o n E n e r g y ( e V )
Photon Energy (eV)
Fig. 2. Micro-PL spectra of a single QD located in a 2-µmdiameter
mesa at two different powers. Inset: Zoom over a
single peak of the spectrum with resolution limited FWHM.
Figure 1 shows the ensemble PL of our sample before and after the RTA process at 775º C
during 5 minutes. The optimization of the growth procedure explained in [1] gives rise to a narrow
emission band centered at 1.296 eV with FWHM=14.6 meV. A blue shift of the PL band and a
reduction of the decay time are clearly observed after the RTA without noticeable change of the
FWHM. T study o the same single QDs before and after the RTA process we have defined arrays of 2-
µm-wide mesa structures on the sample surface (inset Fig. 2). Before the RTA, the s-shell emission is
dominated by multicharged exciton complexes and shows narrow emission linewidths (

TuP23
Ga blocking effect for GaN growth with NH 3
B. Damilano 1,* , A. Kahouli 1,2 , J. Brault 1 , D. Lefebvre 1 , and J. Massies 1
1
Centre de Recherche sur l’Hétéro-Epitaxie et ses Applications, Centre National de la Recherche Scientifique
Valbonne 06560, France
2 Université de Nice Sophia Antipolis, Parc Valrose, 28 avenue Valrose, 06108 Nice cedex 2, France
Controlling the surface stoichiometry during GaN growth is a very important issue since it controls the
surface morphology and material quality. High quality GaN grown with NH 3 is most generally realized
at 800°C with large V/III ratio. Very few works regarding the growth of GaN with ammonia at lower
temperatures and with near-stoichiometry or Ga-rich conditions have been conducted. It has been
shown in previous works [1,2] that for GaN growth with ammonia at temperatures lower than 750°C,
the GaN growth rate as a function of the Ga flux (for a fixed ammonia flux) increases up to a critical
value after which the growth rate starts to decrease.
To get new insights in this peculiar behavior, we have studied the GaN growth on a GaN surface
preliminary exposed to a known amount of Ga. The evolution of the presence of metal (or not) at the
surface has been investigated by recording the (apparent) temperature variation of the sample. The
temperature of the substrate used for these experiments is measured by a narrow-band 1 µm IR
pyrometer. If a metallization of the surface occurs, the transmitted IR signal is modified.
The experiments presented in Fig. 1a) have been conducted at a temperature of 680°C with a Ga flux
and NH 3 fluxes corresponding to a GaN growth rate of R GaN =0.15 ML/s and a V/III ratio of 6. In this
figure are reported the curves corresponding to the relative variation of temperature as a function of
time after the exposition of the GaN surface to Ga during t1 and to Ga+NH 3 during t2 (t1+t2=50sec.).
When no Ga is deposited before Ga+NH 3 exposition, there is only a small change in the observed
temperature variation (curve 1 of Fig. 1a). When only Ga is deposited (curve 4 of Fig. 1a), there is a
decrease of the apparent temperature due to Ga accumulation at the surface, the time needed to recover
a Ga-free surface (Ga desorption time) is indicated by an arrow. When Ga is deposited followed by
Ga+NH3 exposition, intermediate cases are observed (curves 2 and 3 of Fig. 1a). The results are
summarized in Fig. 1b which shows the Ga desorption time as a function of the initial Ga coverage
(t1*R GaN ). The striking feature is that there is a very sharp transition at a Ga coverage of ~2.5 monolayers
above which Ga accumulation is observed. Above this coverage, GaN growth rate is strongly
limited by the blocking of available sites for ammonia adsorption (and cracking) by Ga atoms. It is
worth to note that this value corresponds to the Ga quantity needed to form the stable Ga bilayer
characteristics of Ga-rich GaN (0001) polar surfaces. Implications on different MBE growth strategies
for GaN and alloys using NH 3 will be discussed.
ΔT/T (%)
0
-5
0
-5
0
-5
0
(1)
(2)
(3)
(4)
-5
0 50 100 150 200 250 300
Time (sec.)
Ga desorption time (sec.)
200
150
100
50
0
GaN growth
Ga accumulation
0 1 2 3 4 5 6 7 8
Ga coverage (ML)
(b)
(a)
Fig 1: (a) Relative temperature variation due to Ga deposition during t1 sec. and subsequent Ga+NH3 exposition during t2
sec. on a GaN surface. For all the experiments t1+t2=50 sec., after these 50 seconds the sample stays in vacuum conditions.
For curves 1) to 4) the t1 times are 0, 15, 30, and 50, respectively. The dashed line indicates ΔT/T=0. The time
corresponding to the complete Ga desorption is indicated by an arrow. (b) Ga desorption time as a function of Ga coverage
extracted from the data of Fig. 1a.
[1] R. Held, D.E. Crawford, A.M. Johnston, A.M. Dabiran and P.I. Cohen, Surface Rev. Lett. 5, 913 (1998).
[2] A. Kawaharazuka, T. Yoshizaki, T. Hiratsuka, and Y. Horikoshi, Phys. Status Solidi C 7, 342 (2010).
__________________________
* Contact: bd@crhea.cnrs.fr

TuP24
Use of RHEED to optimize atomic layering of complex oxides
B. A. Davidson 1,* , A. Yu. Petrov 1 and S. Nannarone 2
1
CNR-IOM TASC National Laboratory, Area Science Park-Basovizza, 34149 Trieste Italy
2
Univ. di Modena e Reggio-Emilia, Dipartimento di Ingegneria dei Materiali, Modena, Italy
We have developed a new technique based on RHEED rocking curves that allows us to
quantitatively determine in situ the terminating layer of (001) perovskites ABO 3 (pure AO,
pure BO 2 or the ratio of mixed termination) at any point during the deposition of atomicallyflat
manganite or titanate films. Such real-time, quantitative knowledge is essential for
improving control of the stacking sequence or doping profile when growing interfaces in
which the electronic properties depend sensitively on the atomic structure. Using the reactive
molecular beam epitaxy (MBE), in which each cation is supplied independently at high
precision, this method allows us to identify the film (or substrate) termination and adjust it to
pure AO or BO 2 termination prior to the deposition of the heterointerface. This complements
our recent work to optimize the growth of manganite films by MBE [1]. Such an approach is
also useful, for example, to improve the sharpness of delta-doped layers or the accuracy of
short period superlattices or staircase doping profiles.
[1] B. A. Davidson, A. Yu. Petrov, A. Verna, X. Torrelles, M. Pedio, A. Cossaro, A. Morgante and S. Nannarone, submitted
to APL (2010).
__________________________
* Contact: davidson@tasc.infm.it

TuP25
II-VI nanostructures, with type-II band alignment, for photovoltaics
R. André 1,* , E. Bellet-Amalric 2 , J. Bleuse 2 , C. Bougerol 1 , M. Den Hertog 1 ,
L. Gérard 1 , H. Mariette 1
CEA-CNRS group "Nanophysique et Semiconducteurs"
1
Institut NEEL-CNRS, BP166,38042 Grenoble Cedex 9, France
2 CEA-Grenoble, INAC/SP2M, 17 avenue des Martyrs, 38054 Grenoble Cedex 9, France
In thin film solar cells, based on direct bandgap semiconductors (GaAs, CdTe…) the optical
absorption is very large as compared to the case of the indirect silicon. This is obviously of major
advantage for an efficient sun light collection and it has a direct positive impact on material savings
which are very important when using rare elements. However, a direct bandgap is also highly
favorable for electron-hole radiative recombination, shortening of lifetime and of diffusion length.
Consequently, to fully benefit from those materials as absorbers for photovoltaic (PV), it is necessary
to get rid of their drawback by designing structures which separate spatially electrons and holes to
limit radiative losses.
For that purpose, we studied nanostructures based on the pair of materials CdSe/ZnTe. The
reason is that the interface between those two semiconductors exhibits the so-called type-II, or
staggered, band alignment: the maximum energy of the valence band and the minimum energy of the
conduction band are located at opposite sides of the interface. Excitons photo-generated in the vicinity
of such an interface are then spontaneously dissociated. We have focused specifically on CdSe/ZnTe
because the CdSe bandgap (1.7eV) is adapted to the solar spectrum and because of their very low
lattice mismatch. Nevertheless CdSe/ZnTe is not the only option: a type-II configuration is generally
achieved with II-VI heterostructures when modulating element VI and preliminary results on the
ZnO/CdSe interface will also be discussed.
(a)
5 nm
Fig 1: CdSe/ZnTe superlattice, (a) HR-TEM image, (b) slow PL decay for the interface optical transition at 943 nm (10K).
We have grown, by MBE, two kinds of samples: simple CdSe/ZnTe interfaces or superlattices
(Fig 1(a)). We have been working on tellurium surface segregation and on the related lack of
abruptness of composition at the interface. The formation of alloys may be an issue regarding the
optimization of the band structure for PV because the bandgap of the alloys can be well below the
bandgap of any of their binary components. The structural properties of the samples were
characterized by High Resolution X-ray Diffraction and High-resolution transmission electron
microscopy (HR-TEM).
Regarding optical properties, we have measured, by time resolved photoluminescence (Fig.
1(b)), the efficiency of the electron-hole separation in type-II superlattices. The measured decay time
is above 100 ns for the interface optical transition, i.e. 3 orders of magnitude slower than the typical
PL decay time for the constitutive materials taken separately. This is a direct consequent of the weak
overlap of the electron and hole wavefunctions. We have also studied the effect of the CdSe absorber
thickness which has to be thinner than the diffusion length to fully benefit from the type-II effect, but
kept thick enough to fully absorb light.
__________________________
* Contact: regis.andre@grenoble.cnrs.fr
(b)

TuP26
Enhanced intermixing in Ge nano-prisms on groove patterned Si
(1 1 10) substrates
G. Chen, 1* G. Vastola, 2 J. J. Zhang, 1 B. Sanduijav, 1 G. Springholz, 1 W. Jantsch, 1 F.
Schäffler 1
1 Institute of Semiconductor and Solid State Physics, Johannes Kepler University, Altenbergerstr. 69, A4040 Linz,
Austria
2 Institute of High Performance Computing, 1 Fusionopolis Way, #16-16 Connexis, 138632, Singapore
In the last decade, research on heteroepitaxy of Ge or Si1-xGex on vicinal Si surfaces has
attracted considerable interest because it provides a model system for both growth instabilities
and self-organized surface texturing on a nanoscale. 1, 2 Especially on Si (1 1 10) substrate,
prominent ripple structures could be observed, which are essentially prisms of triangular cross
section that are bounded by two adjacent {105} facets. 3
In this talk, the morphological evolution of {105}-bounded SiGe nano-ripples on
groove-patterned Si (1 1 10) substrates is observed with atomic force microscope for varying
groove widths as shown in Figure 1.
Figure 1 (a)-(d): 2 × 2 µm 2 Laplacian filtered AFM images after 4 ML of Ge deposited on Si (11 10) substrates with various
patterning. (a): flat substrate; (b)-(d): groove-patterned substrates with different groove bottom width L. Inset in (b): 3D
schematics.
The Ge concentration evolution has also been studied with selective chemical etching technique
as shown in Figure 2. 4
Enhanced Si-Ge intermixing between the nano-ripples and the groove sidewalls is interpreted
as the driving force for the observed increase of the ripple volume with decreasing groove width,
and for the reduction of the total number of ripples. Finite element simulations reveal that the
enhanced intermixing arises from the minimization of the total energy density of the ripples.
Our experiments and modeling suggest a direct route for controlling the composition of the
nano-ripples.
[ 1 ] J. Tersoff, Y. H. Phang, Zhenyu Zhang, and M. G. Lagally, Phys. Rev. Lett. 75, 2730 (1995).
[ 2 ] C. Teichert, J. C. Bean, and M. G. Lagally, Appl. Phys. A: Mater. Sci. Process. 67, 675 (1998).
[ 3 ] L. Persichetti, A. Sgarlata, M. Fanfoni, and A. Balzarotti, Phys. Rev. Lett. 104, 036104 (2010).
[ 4 ] A. Rastelli, M. Stoffel, A. Malachias, T. Merdzhanova, G. Katsaros, K. Kern, T. H. Metzger, and O. G. Schmidt, Nano Lett. 8,
1404 (2008).

TuP26
Figure 2 (a)-(b): 3D view of a broad nano-ripple before and after NHH etching for 120 minutes. Inset in (a) and (b):
Laplacian filtered AFM images; (c): a set of profiles corresponding to different etching times measured across the line
C in (a); (d): calculated Ge concentration map based on the etching profiles in (c);
________________________
* Contact: gang.chen@jku.at

TuP27
Influence of nitrogen flux on InGaN growth by PAMBE
H. Turski 1, * , M Siekacz 1, 2 , M. Sawicka 1, 2 , G. Cywiński 1 , M. Kryśko 1 , S. Grzanka 1 ,
I. Grzegory 1 , S. Porowski 1 , Z. Wasilewski 3 and C. Skierbiszewski 1, 2
1
Institute of High Pressure Physics, Polish Academy of Sciences, 01-142 Warszawa, Poland
2
TopGaN Ltd, ul Sokolowska 29/37, 01-142 Warszawa, Poland
3
Institute for Microstructural Sciences, National Research Council, Ottawa, Canada
The role of InGaN as a semiconductor for optoelectronic devices is known to be essential.
Main reason for this is a wide range of accessible band gaps, especially within the visible light
spectrum. That is why so much effort is devoted to develop light emitting diodes of various colors,
green and violet lasers, detectors etc. using this compound. However, growth of high quality InGaN is
very difficult. Large difference between optimum growth temperatures for InN and GaN strongly
limits amount of indium which can be incorporated into the structure. But by changing growth
parameters like gallium and nitrogen fluxes one can influence both InGaN content and surface
morphology.
It was already demonstrated that high quality InGaN structures can be grown on c-plane
(0001) GaN in excess of In flux by Plasma Assisted Molecular Beam Epitaxy (PAMBE). It is also
known that for the growth of InGaN in indium-rich conditions by PAMBE, for constant growth
temperature, Ga/N ratio limits indium content – e. g. for Ga > N only GaN layers are grown.
In this work we studied role of nitrogen flux in InGaN growth. Experiments were done using
different substrates: (a) bulk c-plane GaN substrates with miscut angle (Θ) 0.5° and threading
dislocation density (TDD) of 10 6 cm -2 , grown by Hydride Vapor Phase Epitaxy (HVPE), and (b)
GaN/Al 2 O 3 templates of Θ equal 0.6° with TDD of 10 9 cm -2 . To determine the composition of the
grown samples X-ray diffraction and photoluminescence measurements were done. Surface
morphology of grown samples has been investigated by atomic force microscope.
We found that In content in samples grown using different active nitrogen fluxes differs (see
FIG. 1.) despite that other parameters like the growth temperature, substrate miscut and Ga flux were
kept constant. This effect is caused by the fact that decomposition of In-N fraction is similar for the
same growth temperature. Thus, for greater N flux, rate at which N atoms arrive to atomic kinks is
greater and the InN bond has less time to decompose what implies higher concentration of indium in
the grown layers.
Data obtained from the growths of InGaN with different N and Ga fluxes are presented. It is
shown that for a given growth temperature by increasing nitrogen flux from 4.4∙10 14 [atom/cm 2 ] to 10 15
[atom/cm 2 ] InGaN content can be increased from approximately 14% to almost 20% what is essential
for long wavelength emitters. Surface morphology for different samples is compared. Growth model
describing indium content is discussed. Theoretical results are compared with the experimental data.
Acknowledgements: This work was supported partially by the Polish Ministry of Science and Higher Education Grant No IT
13426 and the European Union within European Regional Development Fund, through grant Innovative Economy
(POIG.01.01.02-00-008/08).
__________________________
*
Contact: Henryk Turski, Phone +48 22 8760352, Fax: +48 22 8760314, email: henryk@unipress.waw.pl

TuP27
20
Ga = 4.9 [nm/min]
Ga = 1.3 [nm/min]
In content [%]
18
16
14
T G
=650 o C
3 4 5 6 7 8 9 10 11 12 13 14
N flux [nm/min]
FIG. 1.
Indium content in function of N flux for two series of samples with the same miscut that were grown at the same temperature.

TuP28
Growth of metal Co/Ag superlattices on MgO(001):
microstructure and magnetic characterization.
Ana Ruiz 1,* , Enrique Navarro 1 , María Alonso 1 , Pilar Ferrer 1,2 ,
Daniel Margineda 1 , F. Javier Palomares 1 , Federico Cebollada 3 , Jesús Mª González 1,4
1
Instituto de Ciencia de Materiales de Madrid-CSIC. Cantoblanco. E28049-Madrid. Spain.
2 SpLine-ESRF. 6, Rue Jules Horowitz. BP 220. F38043 Grenoble Cedex 09. France.
3 Depto. Física Aplicada a las Tecnologías de la Información, Universidad Politécnica. E-28031Madrid. Spain.
4 Unidad Asociada ICMM-CSIC/ IMA-UCM. Apdo. correos 155. E-28230 Las Rozas (Madrid). Spain.
Previous works in the literature have shown the singular and still controversial
behavior of the Co/Ag system as compared to other metallic multilayers. In particular,
perpendicular magnetic anisotropy and oscillations of the magneto resistance and exchange
coupling have only been reported for very thin Co layers (

TuP29
Quantum dot formation from sub-critical InAs layers grown
on metamorphic InGaAs
P. Frigeri 1,* , G. Trevisi 1 and L. Seravalli 1
1
IMEM-CNR Institute, Parco Area delle Scienze 37/A, I-43124 Parma , Italy
Low-density Quantum Dot (QD) structures are currently the object of intensive research devoted to
develop novel nanophotonic devices for quantum communication and computing. In order to tune single-
QD emission at telecom wavelengths (1.3 - 1.55 !m), the growth of InAs/InGaAs metamorphic QD
structures on GaAs substrates, has been successfully proposed [1, 2]. InAs QDs grown on InGaAs show
significant morphological differences with respect to more intensively studied InAs/GaAs system. Since
quantum confinement effects in QD nanostructures are strongly dependent on island shape and island-size
distribution and uniformity, a deeper understanding of QD formation process in metamorphic structures is
essential to improve the prediction and control of their light-emission properties.
Here, we focus on the study of self-aggregation of low density InAs QDs on metamorphic InGaAs buffer
by using Atomic Force Microscopy (AFM) and Photoluminescence (PL) techniques.
InAs QD layers were grown by molecular beam epitaxy on underlying structures consisting of: i) a 100
nm-thick GaAs layer, ii) a 500 nm-thick InxGa1-xAs (x = 0, 0.15, 0.30) and iii) a 5 nm-thick GaAs layer.
Ensemble and single QD optical properties were studied on identical structures capped with 20 nm-thick
In x Ga 1-x As.
The structure design and growth parameters were chosen to enable the study of the self-assembly
mechanism considering the following two concomitant conditions. First, InAs QDs are formed during the
post-growth annealing of an InAs layer thinner than the critical thickness for 2D-3D transition [3].
Avoiding the effects due to incoming In atoms, it is possible to highlight the nucleation mechanisms
dependent on composition-related properties of InGaAs metamorphic layers, such as strain status and
surface corrugation. Second, the insertion of
a thin GaAs layer before InAs deposition, by
reducing the role of the different growth
front-Indium populations associated with the
different buffer compositions, allows to
investigate the effects on QD formation
mainly due to the surface lattice parameter
imposed by the metamorphic InGaAs layer.
We demonstrate that, by optimizing the
values of sub-critical InAs coverage and
Fig 1: 5x5 !m 2 AFM image of QDs formed by postgrowth annealing
(100 s) of a 1.5 ML of InAs grown on metamorphic In 015 Ga 0.85 As.
post-growth annealing time, an accurate
control on island morphology and very low
QD density, down to 10 8 cm -2 (Fig 1), can be
achieved. Moreover, micro-PL experiments performed on InAs/In 015 Ga 0.85 As structures reveal an efficient
single-QD emission [4].
Finally, we discuss the challenges arising from the combined use of metamorphic and sub-critical InAs
deposition approaches to drive the positioning of QDs on the surface, an essential requisite for quantum
information applications.
[1] E. S. Semenova et al, J. Appl. Phys., 103, 103533 (2008).
[2] L. Seravalli, G. Trevisi, P. Frigeri and C. Bocchi, Journal of Physics:Conference Series, 245, 012074 (2010)
[3] H.Z. Song, T. Usuki, Y. Nakata, N. Yokoyama, H. Sasakura and S. Muto, Phys. Rev. B, 73, 115327 (2006)
[4] J. P. Martinez Pastor, to be published
* Contact: pfrigeri@imem.cnr.it

TuP30
Reversible Nanofacetting and 1D Ripple Formation of
Ge on High-Indexed Si (11 10) Substrat
G. Springholz*, B. Sanduijav and D. Matei
Institute for Semiconductor Physics, Johannes Kepler University, Altenbergerstr. 69, A-4040 Linz, Austria
Silicon-germanium has been an intensely studied model system for the growth of self-assembled
quantum dots by the Stranski-Krastanow mode. A prominent feature of these dots is their highly
facetted pyramidal or dome-like shape [1] that is governed by the formation of energetically favored
side facets. The formation of the different dot shapes not only depends on the Ge coverage and growth
condition, but also on the substrate orientation [2]. The (11 10) Si substrate orientation is special in this
respect because of its particular relationship to the low energy {105} facets of compressively strained
Ge pyramids, in which case the intersections of these facets are parallel to the (1110) surface. In
addition, local (11 10) Si substrate facets also play a major role in site-controlled growth of Ge islands
on pit-patterned or stripe-patterned Si substrate templates [3,4].
In the present work, Ge growth on high-indexed (11 10) substrates was studied systematically using in
situ scanning tunneling microscopy. The experiments were performed in a multi-chamber MBE/STM
system, allowing sequential growth and imaging of the epitaxial surface structure formed after each
growth step [4]. The results demonstrate that the (1110) growth properties radically differ from those
on the usual (001) Si substrates. As shown by the sequence of STM images depicted in Fig. 1, at
certain critical coverage of ~ 4 monolayers (ML), a highly stable quasi-periodic 1D ripple structure is
formed perpendicular to the dimer direction. As demonstrated by Fig. 1(a) below the critical coverage
of 4 ML only a random surface roughening takes place, but within the next fractional monolayer an
abrupt transformation to well defined {105} facetted ripples occurs (see lower panel of Fig. 1). These
ripples completely cover the whole (1110) substrate surface, in contrast to the usual isolated Ge islands
formed by the Stranski-Krastanow growth mode.
A quantitative analysis shows that in the ripple formation process, the initial 2D Ge wetting layer is
completely consumed, i.e., no wetting layer remains underneath the ripples. Thus, the ripples represent
a novel pathway for lowering the free energy of the system. Moreover, the ripples show a well defined
and unique width and height with a lateral periodicity of ~20 nm with a rather narrow size dispersion.
To assess the thermal stability, post growth annealing experiments were performed and compared with
annealing of Ge islands grown on Si (001) under the same growth conditions. During annealing, an
elongation of ripples is observed but only a small change in ripple size and period occurs even after
long term annealing, whereas for the Ge islands on (001) the usual Ostwald ripening occurs during
which the island density continuously decreases. Most strikingly, during thermal cycling of the Ge
layers on Si (1110), a reversible transition from the rippled surface to a flat surface is found to occur
when the annealing temperature is raised above a critical temperature of 600°C. This process is
reversed when the temperature is lowered again. This is demonstrated by the RHEED patterns shown
on the left hand side of Fig. 2, where a 5 ML Ge layer on Si (1110) was slowly heated from 500 to
620°C and then cooled back to 500°C. The intensity evolution of one of the RHEED diffraction spots
characteristic for the rippled Ge surface as a function of the sample temperature is shown on the right
hand side of Fig. 2 during multiple heating and cooling cycles. Both factors clearly demonstrate, that
the ripples represent an equilibrium structure and their formation closely resembles a thermodynamic
phase transition.
[1] see M. Brehm, et al., Nanoscale Research Letters (in print) and references therein.
[2] J.T. Robinson, A. Rastelli, O. Schmidt and O.D. Dubon, Nanotechnology 20, 085708 (2009).
[3] G. Chen, H. Lichtenberger, G. Bauer, W. Jantsch, F. Schaeffler, Phys. Rev. B 74, 035302 (2006).
[4] B. Sanduijav, D.G. Matei, G. Chen, G. Springholz, Phys. Rev. B 80, 125329 (2009).
__________________________
* Contact: gunther.springholz@jku.at

TuP30
Fig. 1: Scanning tunneling microscopy images of Ge on Si (11 10) substrates at increasing Ge coverage from
1.8 to 5.3 ML grown by MBE at 550°C. The surface orientation maps of the STM images are shown as inserts.
Fig. 2: Reversible ripple formation of 5 ML Ge on Si (11 10) substrates revealed by in situ RHEED investigations.
Left: Sequence of RHEED patterns recorded during heating and cooling of the surface from 500°C to
620°C and back to 510°C. Right: Intensity of the 3D ripple spot indicated by the dashed rectangle in the
RHEED patterns of the left hand side measured as a function of substrate temperature during multiple heating
and cooling cycles of the static 5 ML Ge surface from a temperature of 450°C to 620°C.

TuP31
Homoepitaxy and nitrogen doping of non polar ( 1010)
ZnO films
D. Tainoff 1 , J.-M. Chauveau 1,2 , C. Deparis 1 , B. Vinter 1,2 , M. AlKhalfioui 1,2 , M.
Teisseire 1 , Christian Morhain 1 .
1 CNRS-CRHEA, Av. Bernard Grégory, F- 06560 Valbonne Sophia Antipolis, France
2
University of Nice Sophia Antipolis, Parc Valrose, F-06102 Nice Cedex 2, France
In spite of many studies, p-type doping of ZnO and its related alloys is still a blocking point for the
development of ZnO based optoelectronic devices. Among many candidates (As, P, N, C …) studied
to achieve p type doping of ZnO, the more promising results have been obtained using nitrogen by
Tsukasaki et al. 1 In this case the free hole concentration at room temperature is limited to 10 16 cm -3 .
However high p type doping level could not be achieve possibly because most of these studies
were based on ZnO grown on foreign substrates (Sapphire, Si, SCAM, LiAlO 3… ). Indeed, the
heteroepitaxy of ZnO generates a high density of impurities and structural defects leading to local
inhomogeneities of the dopants and parallel conduction channels. Therefore, reliable electrical
characterizations are not straightforward on heteroepitaxial layers. The use of ZnO bulk substrates
should circumvent these limitations. In addition, it has already been shown that the orientation of the
substrate plays an important role in the intentional and unintentional incorporation of dopants,
including nitrogen which seems to be facilitate in case less polar polar orientation 4 .
In this presentation we report on the interplays between growth parameters, structural, optical and
electrical properties in non intentionally doped (n.i.d.) and nitrogen doped ZnO thin films grown on
non polar (1010) oriented (m plane) ZnO substrates by MBE. We show that the N incorporation and
the growth temperature (T Gr ) do not affect significantly the growth process on a large growth window.
Indeed, RHEED measurements exhibit streaky patterns in both doped and undoped films for T Gr from
370°C to 550°C. The RMS roughness measured by AFM for doped and undoped film is below to 10 Å
all over the temperature range studied here. Therefore, the surface morphology of ZnO layers is
comparable, which allows to unambiguously study the dopant incorporation as a function of the T Gr .
The incorporation of residual impurities and nitrogen has been studied by Secondary Ion Mass
Spectroscopy (SIMS). The donor type residual impurities concentration is dominated by Al
(~10 15 cm -3 ) and does not depend on T Gr . This low concentration of donor type impurities is confirmed
by C (V) measurements (figure 1.a) which show a record value of residual doping of about 10 -14 cm -3
for the (1010) homoepitaxial films.
Nitrogen was activated in an rf-plasma cell. SIMS measurements show that the concentration
strongly depends on the growth temperature (T Gr ) and can be varied from 3.10 18 cm -3 to 10 20 cm -3 . Low
temperature photoluminescence spectra taken from ZnO:N layers are shown in figure 1b). A broad
donor acceptor pair (DAP) emission is systematically observed around 3.24 eV for nitrogen doped
films which is an unambiguous optical signature of shallow acceptor levels. The DAP position slightly
shifts towards higher energies with the increase of nitrogen concentration owing to the decrease of the
distance between the donor and the acceptor. Electrical measurements on nitrogen doped films show
highly compensated n-type films. Since the residual donor concentration, the origin of this
compensation is assigned to the incorporation of N 2 and/or donor type complexes related to nitrogen.
This work was partially funded through the French Carnot Program on Solid State Lighting, under a
collaboration scheme with CEA/LETI
[1] A. Tsukazaki et al. Nat. Mat. 4, 42 (2005)
[2] J.L. Lyons et al. Appl. Phys. Lett. 95, 252105 (2009)
[3] F. Gallino et al. J. Mat. Chem. 20, 689 (2010)
[4] P. Fons et al. Phys. Rev. Lett. 96, 045504 (2006)
_______________________
* Contact: dt@crhea.cnrs.fr

Concentration [cm -3 ]
Intensity (a.u.)
TuP31
a)
1E19
1E18
1E17
1E16
1E15
1E14
Li
EPILAYER
Al
Ga
SUBSTRATE
N D
-N A
0,0 0,2 0,4 0,6 0,8 1,0 1,2 1,4
Depth (µm)
ZnO n.i.d
[N]=1 10 19 cm -3
[N]=7 10 18 cm -3
DAP-2LO
DAP-1LO
DAP
Donors
Bound Excitons
FX
3,0 3,1 3,2 3,3 3,4
b)
Energy (eV)
Figure 1: a) Impurities concentrations determined by SIMS in n.i.d ZnO (1010) homoepitaxial layer (Al, Ga, Li). The
SIMS profile is consistent with the C (V) profile (black square) for both the layer and the substrate. b) Photoluminescence
spectra from nitrogen doped ZnO thin films. The nitrogen doped spectra are normalized on their maximum intensities. The
DAP shift related to the lowering of the donor acceptor distance appears clearly.

TuP32
Structural characterization of MBE grown GaP/Si nanolayers
W. Guo 1 , G. Elias 2 , A. Létoublon 1 *, C. Cornet 1 , A. Ponchet 2 , A. Bondi 1 , T. Rohel 1 , N.
Bertru 1 , C. Robert 1 , T. Nguyen Thanh 1 , O. Durand 1 , J.S. Micha 3 and A. Le Corre 1
1 Université Européenne de Bretagne, France
INSA, FOTON, UMR 6082, F-35708 RENNES
2 CEMES, UPR CNRS 8011,F-31055 Toulouse, France
3 UMR SPrAM 5819 CNRS-CEA-UJF, INAC, F-38054 Grenoble, France
In the context of monolithic integration of photonics on silicon, growth of GaP (III-V
semiconductor) directly deposited on Si has been proposed to overcome the problems of lattice
mismatch. This opens the route for direct bandgap growth on GaP-Si pseudo-substrate. [1],[2]
However, long-term stable device performance implies reproducible achievement of defect-free
interfaces between III-V and Si. Different defects can be generated at the GaP/Si interface: among
them antiphase domains (APD) and microtwins (MT) (see figure 1) are quite difficult to avoid. Their
density and their emergence at the GaP surface must be limited for the subsequent growth of an active
area. In this paper, X-ray diffraction (XRD) on laboratory setups and on synchrotron is combined with
HRTEM and AFM analysis for characterization of defects in GaP thin films grown onto Si(001)
misoriented substrates.
The 20 nm GaP layer has been grown using Migration Enhanced Epitaxy (MEE) growth mode
in order to enhance the smoothing of the growth front, on a (001) oriented Si 4°-off substrate to favour
double Si steps. [1] GaP is deposited after a chemical (modified conventional RCA) preparation and
thermal (10 min at 900 °C) treatment of the Si surface. The growth temperature (Tg) of the substrate
was set at 350°C. This temperature ensured a low roughness of the 20 nm GaP surface measured by
AFM (RMS of 1.2 nm) as compared to Tg = 450 °C (RMS = 2.6 nm). Synchrotron XRD has been
employed in grazing incidence condition to enhance the contribution of the GaP layer. The
contribution of emerging APD with a characteristic broadening of the weak reflections is shown on
figure 2a around the GaP(2 -2 2). [3] This coincides with results obtained on a laboratory setup. [4]
Weak but sharp diffuse streaks, unobserved on the laboratory setups, also show up around this
reflection along three fold axes. Broad reflections are found at positions in agreement with the
presence of microtwins (MT) which correspond to a 180° rotation of the GaP main phase around a
[111] axis [5]. The 4 possible MT variants have been identified on similar scans. The MT average
thickness (T on fig. 1b) has been also evaluated as the inverse of the integral breadth of the broad
peak. For fig. 2b a correlation correlation length of about 6 nm is found across the MT reflection at
(1.66, -2.33, 1.66). High resolution transmission electron microscope (HRTEM) has also been
performed on a sample grown in the same conditions. As shown on figure 3a, large step bunches,
compared to the targeted double Si steps, are observed at the Si surface. Numerous defects are also
evidenced and most of them originate from the GaP/Si interface. Among them, as shown on figure 3b,
MT and Stacking Faults (SF) can be clearly identified in the GaP layer.
As a conclusion, by decreasing the Tg from 450 to 350°C, the surface roughness has been
successfully lowered to 1.2nm. However, a combined analysis XRD, and HRTEM revealed structural
defects in too high density for subsequent photonic applications. These defects are probably related to
the presence of impurities at the Si surface. Thanks to a buffer layer growth, a higher Si surface quality
should allow the growth of high structural quality GaP. [6]
[1] H. Yonezu, Y. Furukawa, A. Wakahara, J. Crystal Growth, 310 4757 (2008).
[2] W. Guo, A. Bondi, C. Cornet, et al., Phys. Status Solidi (c), 6 2207 (2009).
[3] D. A. Neuman, H. Zabel, R. Fischer, H. Morkoç, J. Appl. Phys.,. 61 1023 (1987).
[4] A. Létoublon, et al., J. Crys. Growth, doi:10.1016/j.jcrysgro.2010.10.137 (2010).
[5] K. Hiruma et al., Journ. of Appl. Phys., 77 447 (1995).
[6] T. J. Grassman, M. R. Brenner, et al., Appl. Phys. Lett., 94, 232106 (2009).
[7] I. Németh, Ph. D. Thesis, Philips Univ. Marburg, Germany (2008).
__________________________
* Contact: antoine.letoublon@insa-rennes.fr

TuP32
Fig 1: Typical defects in GaP/Si nanolayers with emerging APD (a) after Nemeth et al. [7] and a MT (b).
Fig. 2 Transverse RSM (a) around the (2 -2 2) reflection showing MT (with streaks along the [1 1 1] type directions)
and E-APD (with diffuse scattering enhanced along the [1 1 0] direction). A scan along the [1 1 1] streak (b) exhibits
a clear MT contribution at about (1.66, -2.33, 1.66).
Figure 3: HRTEM images showing step bunching of the Si surface (a) and several planar defects originating from the
GaP/Si interface (b).

TuP33
An Auger Electron Analyzer System for In situ MBE Stoichiometry
Control
W. Laws Calley 1 ,* Phillipe Staib 2 , Jonathan Lowder 1 , and W. Alan Doolittle 1
1 School of Electrical and Computer Engineering, Georgia Institute of Technology, Atlanta, GA 30332, USA
2 Staib Instruments, Williamsburg, VA 23185, USA
Auger Electron Spectroscopy (AES) analysis is an excellent surface sensitive technique for
thin film analysis and is sensitive to almost all elements [1]. Not only can AES help determine the
species present at the surface, but AES can also yield information about the chemical bonding of the
film [1]. AES also exhibits higher spatial resolution than alternatives like X-ray Photoelectron
Spectroscopy [1]. However, this analysis tool has historically been an ex situ technique with a few
noted exceptions [2]. Recently, Staib Instruments has developed a MBE-specific Auger electron
detection head, the Staib In situ Auger Probe (SIAP), which can be used in situ with a large enough
working distance, tested up 82 mm, so as to not interfere with a growing MBE film and leverages their
existing RHEED gun as the e-beam source.
This in-situ technique works by positioning the auger head normal to the growth surface. The
analyzer head is inserted as close to the sample as possible in order to maximize the capture of the
Auger electrons without interfering with the source fluxes. In the present case this standoff distance is
roughly 5-8 cm. Using the normal-emission detector geometry, the high energy electrons that
contribute to the RHEED image are avoided and the bulk of the electrons captured by the detector are
comprised of secondary and Auger electrons. A built-in energy filter is electronically scanned to give a
spectrum of electrons versus energy as shown in figures 1 - 3. Sensitivity can be improved by
implementing a lock-in detection methodology. This also allows the RHEED system to function
normally while Auger analysis is preformed during the growth. Unlike a traditional ex situ AES
system, the design of the in situ system – specifically the longer working distance and lower energy
resolution – targets control of compositional variation (stoichiometry) in situ whereas information on
chemical bonding is harder to determine.
In tests, a silicon wafer was used in a Varian Gen II MBE system with Tb, Dy, and Fe sources.
The initial test showed oxygen and carbon present on the surface of the Si wafer, figure 1. When Fe
flux impinged on the surface of the silicon wafer the deposition of Fe was observed via the SIAP,
figure 2. Further tests were conducted starting with a layer of Dy and depositing part of a monolayer
of Tb in 2 percent increments. Figure 3 shows a clear distinction between bare Dy and 2, 4, 6, 8, and
10 percent of a monolayer coverage of Tb on the Dy layer, demonstrating the Auger probes sensitivity
to heavy elements. Finally, a near normal incident electron gun was used at low energy, 1keV, to test
the Auger probes ability to detect ELLs lines, figure 4.
This in situ Auger system is very promising as a complementary MBE analysis tool to the
existing RHEED system. While RHEED gives information on the crystal structure and morphology of
the growing film it is now possible to gather chemical information about the growing film. This can be
especially useful when working with films without line compositions or where stoichiometry control is
problematic. As the energy resolution is improved, growth of films that have multiple oxidation states
or other similar phase/chemical transitions that are important to monitor during growth might be
attainable through a closed loop control system. This promising technique would also yield a nondestructive
“depth profile” since the data would be taken as the film was grown instead of drilling the
film after growth to perform the analysis. Finally this technique could give information about
transitions between layers in multilayered films grown via MBE.
[1] D. P. Woodruff, T. A. Delchar, Modern Techniques of Surface Science – Second Edition. Cambridge University Press
(1999).
[2] M.L. Watson, J.S.S. Whiting, A. Chambers, and S.M. Thompson, J. Magn. Magn. Mater. 113, 97-104 (1992).
*Contact: laws@gatech.edu

TuP33
Intensity [Arb. Units]
Carbon
Oxygen
Silicon
200 400 600 800 1000 1200 1400 1600 1800
Energy (eV)
Fig. 1 Scan of a Si wafer after loading with oxygen and
carbon on the surface.
Fig. 2 Iron peaks detected on during iron deposition on a
silicon wafer.
Fig. 3 Dysprosium on a Si wafer with different percent
monolayer coverage of Terbium ranging from bare Dy to
10% of a monolayer coverage of Tb.
Fig. 4 ELLs on a Si wafer using a 1keV electron beam.

TuP34
MBE droplet epitaxy of InGaAs/Ga (100): effect of MBE conditions and post-annealings on (Ga,In)
droplet and GaInAs nanostructure morphology and emission
Poonyasiri Boonpeng 1,2,3 , Guy Lacoste 1, 2 , Alexandre Arnoult 1,2 , Hejer Makhloufi 1,2 , Olivier Gauthier-
Lafaye 1,2 , Guilhem Almuneau 1,2 , Somchai Ratanathammaphan 3 , Somsak Panyakeow 3 , Chantal
Fontaine* 1,2
1 CNRS ; LAAS ; 7 avenue du colonel Roche, F-31077 Toulouse, France
2
Université de Toulouse ; UPS, INSA, INP, ISAE ; LAAS ; F-31077 Toulouse, France
3 Semiconductor Device Research Laboratory (Nanotec Center of Excellence),
Department of Electrical Engineering, Faculty of Engineering, Chulalongkorn University, Bangkok
10330, Thailand
In x Ga 1-x As nanostructures on GaAs(001) substrates grown by droplet epitaxy were achieved
using molecular beam Epitaxy (MBE). [1,2] The influence of experimental MBE conditions
on In x Ga 1-x droplet characteristics were investigated, as well as on In x Ga 1-x As quantum
nanostructures after crystallization under As 4 flux. We will discuss how these parameters play
a part upon nanostructures characteristics, as observed by atomic force microscopy (AFM).
Some of these nanostructures, embedded in GaAs, were annealed using rapid thermal
treatment. These were studied by photoluminescence spectroscopy. We will present how
emission evolves with annealing.
*chantal.fontaine@laas.fr

TuP35
Synthesis of AlGaN nanowires by Molecular Beam Epitaxy
A. Pierret 1,2,* , C. Bougerol 1 , B. Gayral 1 , B. Attal-Trétout 3 , A. Loiseau 2 and B.
Daudin 1
1
CEA-CNRS group "Nanophysique et Semiconducteurs", Institut Néel/CNRS-Univ. J. Fourier and CEA
Grenoble, INAC, SP2M, 17 rue des Martyrs, 38 054 Grenoble, France
2
ONERA – Laboratoire d’Etude des Microstructures, UMR 104 ONERA-CNRS, 29 avenue de la Division
Leclerc BP 72, 92322 Châtillon cedex – France
3
ONERA – Département Mesures Physiques, 27 chemin de la Hunière 91761 Palaiseau cedex - France
Aluminium gallium nitride (AlGaN) is a promising ternary nitride compound because of its
potential applications in optoelectronics devices, by tuning their emission wavelength between 350 nm
(GaN bandgap wavelength) and 200 nm (AlN bandgap wavelength). But for applications, the
comprehension of the optical properties of AlGaN alloys as well as a drastic improvement of structural
ones is requested. Along these lines the use of nanowires (NWs) is particularly attractive, due to their
excellent crystallographic quality. Then, the aim of this work is to address the issue of AlN and
AlGaN NWs growth by Molecular Beam Epitaxy (MBE). One peculiarity of this method is that it does
not require the use of catalyst, ensuring a high chemical purity, as demonstrated for GaN NWs, which
has been favorable for the study of intrinsic optical properties. For AlN, the first reported synthesis of
such structure by MBE is very recent [1]. As concerns AlGaN material, bidimensional layers have
been widely studied these last years and still suffer of a high density of crystallographic defects. This
induces a poor optical efficiency of the UV LEDs (typically 3% at the state of the art, at 260 nm). In
order to overcome it, we propose here to develop the growth of AlGaN NWs. To date, only scarce
results concerning the growth of such NWs using different techniques (MOCVD, MBE, hot-wall
chloride vapor epitaxy) have been reported in literature. We present here results of AlGaN NWs
growth by MBE.
Up to now, AlGaN synthesis by MBE has been achieved by self-induced nucleation on
Si(111), as for GaN NWs. It has been shown that besides having NWs, a 2D layer grows between
them, and is thicker by increasing the aluminium flux [2]. Preliminary cathodoluminescence
experiments have shown that this layer has a high aluminium content whereas NWs are almost pure
GaN. This can be explained by a diffusion length shorter for aluminium than for gallium. Moreover
for high aluminium content, the lattice matching between Si(111) and AlN(0001) prevents the
occurrence of the 2D/3D transition, necessary for GaN NWs growth on Si(111). As a consequence we
can expect that for AlGaN with a high Al content, Si is totally wetted and thus a 2D layer is easily
grown.
A first step toward the understanding of AlGaN NWs growth is to quantify the difference in
aluminium and gallium mobility and a possible influence of gallium in the diffusion of aluminium by
surfactant effect. For this purpose, we have grown AlGaN NWs with different nominal Al/Ga ratio
values and temperatures, on a GaN NW basis grown on Si (111). Photoluminescence, EDX and TEM
experiments have been performed to analyze the actual composition of the NWs. It has been found
that the NW Al content is lower than expected, as a clue that Al poorly diffuses along the sidewalls of
the NWs whereas incorporated Ga comes from both impinging flux on top and flux diffusing along the
sidewalls. We also found that the 2D layer connecting the NWs exhibits a larger Al content than
expected from the Al/Ga ratio used, as a further clue that Al diffusion is very reduced compared to that
of Ga.
[1] O. Landre and al., Appl. Phys. Lett., 96, 061912 (2010).
[2] J. Ristic and al., Phys. Stat. Sol. (A), 192, 60 (2002).
__________________________
* Contact: aurelie.pierret@cea.fr

TuP36
Correlation of structural, chemical and optical characterization
of CdSe quantum dots inserted in ZnSe nanowires
M. Elouneg-Jamroz * , M. den Hertog, S. Bounouar, E. Bellet-Amalric, R. André,
Y. Genuist, K. Kheng, J-P Poizat, S. Tatarenko
Nanophysics and Semiconductors Group, INAC and Institut NEEL, CEA/CNRS/University Joseph Fourier,
25 rue des Martyrs, Grenoble, France
The incorporation of CdSe quantum dots (QDs) as a heterostructure segment in ZnSe nanowires
(NWs) will be presented as part of an effort to produce single photon emitters with tunable energy [1].
We synthesize these NWs by MBE with separate effusion sources of Zn, Cd and Se and the growth is
catalyzed by gold which we evaporate onto the sample from a solid source. The growth is initiated on
ZnSe 2D layers epitaxially grown on (100) and (111)B GaAs substrates. For both orientations NWs
follow epitaxial directions, although it is generally less clear in the (100) case. [100], [111] cubic and
[0001] hexagonal NWs are obtained.
For the structural and chemical characterization of the CdSe QDs we rely solely on TEM detection.
We will show results obtained by HRTEM, HAADF-STEM, EFTEM and EDX.
These are exceptional results given these 10nm thin NWs have a very short lifetime under the TEM
beam.
µ-Photoluminescence spectra performed on single NWs show characteristic exciton and biexciton
emission lines. Anti-bunching experiments are performed on a typical HBT µ-photoluminescence
setup.
The crystalline structure, the chemical composition and size of the QD will be discussed for a few
samples. A correlation between QD size and luminescence is observed. Preliminary results show a
blue shift from 2.2 to 2.45eV going from the largest to the smallest QD.
2700
2500
em ission (m eV )
2300
2100
1900
1700
20 30 40 50 60 70 80
size QD (A)
Fig 1: µ-photoluminecence emission energy of the
exciton line in single CdSe QDs vs the QD length
[1] A. Tribu et al. Nano Lett. 8, 4326 (2008)
[2] E. Bellet-Amalric, M. Elouneg-Jamroz,, C. Bougerol, M. Den Hertog, Y. Genuist, S. Bounouar, J.P. Poizat, K. Kheng, R.
André and S. Tatarenko Physica Statu Solidi C 7 (6) 1527 (2010)
__________________________
* Contact: miryam.elouneg-jamroz@cea.fr

TuP36
Fig 2:ZnSe NW grown at 400°C with a
30s CdSe insertion grown on GaAs (111)B
(A) SEM image of the NWs
(B) HAADF STEM image of two NWs side by side, a
bright contrast in the NWs is observed 25 nm from
the catalyst particle, marking the QD insertions.
(C) Zoom of the region marked in (B)
(D) Intensity profile obtained along the dotted rectangle
in (C). The QD is clearly visible and is 3.9 nm long

TuP37
Optical and structural properties of InGaN/GaN nanowires
G. Tourbot 1,2* , C. Bougerol 3 , C. Leclere 4 , B. Gayral 2 , P. Gilet 1 , H. Renevier 4
and B. Daudin 2
1
CEA-LETI, Minatec Campus, 17 Rue des Martyrs 38054 Grenoble Cedex 09
2 CEA-CNRS-UJF group « Nanophysique et Semiconducteurs », CEA, INAC, SP2M, NPSC, 17 rue des Martyrs,
38 054 Grenoble, France
3 CEA-CNRS-UJF group « Nanophysique et Semiconducteurs », Institut Néel/CNRS, 25 rue des Martyrs,
38 042 Grenoble, France
4
Laboratoire des Matériaux et du Génie Physique, Grenoble INP - MINATEC, 3 parvis L. Néel
38016 Grenoble, France
The nanowire structure offers a very interesting way of bypassing the high lattice mismatch
and density of dislocation in 2D InGaN/GaN heterostructures. This opens a way towards the
realization of efficient III-N-based LEDs involving highly mismatched InGaN/GaN in the green and
yellow spectral range, whose performances currently lag behind those of less mismatched blue and
violet LEDs. Green-emitting nanowire-based InGaN/GaN LEDs have been demonstrated, but a great
deal of optimization remains to be done before such objects can match the performances of their 2D
counterparts. In this aim, the understanding and control of the properties (both structural and optical)
of the InGaN active region is crucial.
(a)
(b)
Fig 1: (a) InGaN quantum dots in a GaN nanowire and (b) comparison of their PL emission et 7K and 300K,
normalized by the integration time
We have demonstrated that the growth of InGaN/GaN nanowires by MBE can result in a
spontaneous phase separation into an In-rich core and an In-poor shell [1]. We will present a detailed
structural study on the respective strain states and compositions of the core and shell using X-ray
diffraction and Diffraction Anomalous Fine Structure.
Based on the growth mechanism proposed in [1], the structural properties of InGaN axial
insertions in GaN nanowires (fig. 1a) will be detailed, including the influence of the insertion size on
In incorporation. We will also demonstrate that such structures, offering a spontaneous lateral
confinement of excitons preventing surface recombination, are strong emitters up to room temperature,
even in the green spectral range.
[1] G. Tourbot et al., submitted to Nanotechnology (2010).
__________________________
* Contact: Gabriel.tourbot@cea.fr

TuP38
Position controlled self-catalyzed growth of GaAs nanowires by molecular
beam epitaxy
Andreas Rudolph, Joachim Hubmann, Markus Kargl, Benedikt Bauer, Marcello Soda,
Josef Zweck, Dieter Schuh, Dominique Bougeard and Elisabeth Reiger
Institute for Experimental and Applied Physics, University of Regensburg, Universitätsstr. 31,
D-93053 Regensburg, Germany
andreas.rudolph@physik.uni-regensburg.de
Anna Fontcuberta i Morral
Laboratoire des Matériaux Semiconducteurs, Institut des Matériaux, École polytechnique
fédérale de Lausanne, CH-1015 Lausanne, Switzerland
Nanowires grown in bottom-up processes are regarded as possible building blocks for future
electronic devices. A major challenge on the way of integrating nanowires in conventional
electronic circuits is the control of the diameter and the position. We report on position
controlled GaAs nanowires grown via self-catalyzed and Au catalyzed growth using MBE
[1,2]. For the self-catalyzed growth we use GaAs (111)B wafers covered by a thin SiO 2 layer
as substrate. With E-beam lithography in combination with wet chemical etching
arrays of holes with diameters between 100nm and 400nm and varying interhole distances
between 200 and 2000 nm are defined. The etched holes in the SiO2 layer act as nucleation
sites for nanowire growth. The nanowires are mostly oriented in the [111]B direction,
nanowire growth is restricted to the patterned areas. SEM/TEM characterizations show that
the nanowires have a hexagonal shape with {110} side facets and zinc blende as dominant
crystal structure.
For the Au-catalyzed growth gold disks with diameter between 100nm and 400nm at a fixed
interhole distance of 2000 nm are fabricated using E-beam lithography and Lift-Off
technique. Due to migration of Au, the growth of several NWs per gold disc is observed.
[1] C. Colombo, D. Spirkoska, M. Frimmer, G. Abstreiter, A Fontcuberta i Morral., Phys. Rev. B 77 (2008)
155326.
[2] Benedikt Bauer, Andreas Rudolph, Marcello Soda, Anna Fontcuberta i Morral, Josef Zweck, Dieter Schuh
and Elisabeth Reiger, Nanotechnology 21 (2010) 435601.

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LIST OF PARTICIPANTS
ALEKSANDROVA Anna, Humboldt University (GERMANY)
ALÉN Benito, IMM, CSIC (SPAIN)
ALEXEEV Alexey, SemiTEq JSC (RUSSIA)
ALONSO ALVAREZ Diego, IMM, CSIC (SPAIN)
ANDRE Régis, Institut Néel, CNRS (FRANCE)
ANDREWS Aaron Maxwell, Vienna Univ. of Technology (AUSTRIA)
ARNOULT Alexandre, LAAS, CNRS (FRANCE)
AS Donat, University of Paderborn (GERMANY)
ASAOKA Hidehito, Japan Atomic Energy Agency (JAPAN)
AUBERT Arnaud, ALTEC Equipment (FRANCE)
BADANO Giacomo, CEA-Grenoble, LETI (FRANCE)
BELAHSENE Sofiane, IES, CNRS (FRANCE)
BELLET AMALRIC Edith, CEA-Grenoble, INAC (FRANCE)
BIASIOL Giorgio, Ist. Officina dei Materiali CNR (ITALY)
BICHLER Max, Walter Schottky Institut (GERMANY)
BIERWAGEN Oliver, Paul-Drude-Institut (GERMANY)
BIETTI Sergio, Università di Milano Bicocca (ITALY)
BISPING Dirk, Veeco Instruments Inc. (GERMANY)
BLAIN Vincent, Vinci Technologies (FRANCE)
BOISSIER Guilhem, IES, CNRS (FRANCE)
BONANNI Alberta, Johannes Kepler Univ. (AUSTRIA)
BORISOVA Svetlana, Forschungszentrum Jülich (GERMANY)
BOUCHAIB Pierre, RIBER (FRANCE)
BOUGEARD Dominique, Universität Regensburg (GERMANY)
BOUGEROL Catherine, Institut Néel, CNRS (FRANCE)
BRAULT Julien, CRHEA, CNRS (FRANCE)
BRESNAHAN Rich, Veeco Instruments Inc. (USA)
CALLEJA Enrique, ISOM (SPAIN)
CALLEY Laws, Georgia Inst. of Technology (USA)
CERUTTI Laurent, IES, CNRS (FRANCE)
CHAIX Catherine, RIBER (FRANCE)
CHAKRABARTI Subhananda, ITT-Bombay (INDIA)
CHAUVEAU Jean-Michel, CRHEA, CNRS (FRANCE)
CHETTAOUI Azza, Ecole Centrale de Lyon (FRANCE)
CHEZE Caroline, TopGaN Ltd. (POLAND)
CHO YongJin, Paul-Drude-Institut (GERMANY)
CIBERT Joel, Institut Néel, CNRS (FRANCE)
CLARKE Roy, University of Michigan (USA)
COOMBER Stuart, Wafer Technology Ltd (UK)
CORNET Charles, FOTON - INSA (FRANCE)

COUPPEY Catherine, RIBER (FRANCE)
CURE Yoann, CEA-Grenoble, INAC (FRANCE)
DAMILANO Benjamin, CRHEA, CNRS (FRANCE)
DANSHEYA Anna, RIBER (FRANCE)
DAS Aparna, CEA-Grenoble, INAC (FRANCE)
DAUDIN Bruno, CEA-Grenoble, INAC (FRANCE)
DAVYDOK Anton, Universität Siegen (GERMANY)
DESPLANQUE Ludovic, IEMN, CNRS - Univ..Lille (FRANCE)
DIAT Jean Luc, FLEX-EQUIP. / MEWASA (FRANCE)
DIMAKIS Emmanouil, Paul-Drude-Institut (GERMANY)
DOGAN Pinar, Paul-Drude-Institut (GERMANY)
DONATINI Fabrice, Institut Néel, CNRS (FRANCE)
DOSANJH Jeevan, RTA Instruments, (UK)
DUCRUET Marion, CEA-Grenoble, INAC (FRANCE)
DUPUY Emmanuel, CEA-Grenoble, INAC (FRANCE)
DUSSAUD Jean, Institut Néel, CNRS (FRANCE)
EBEL Lars, University Wuerzburg (GERMANY)
EIBELHUBER Martin, Johannes Kepler Univ. (AUSTRIA)
EICKHOFF Martin, Justus-Liebig-University (GERMANY)
ELOUNEG-JAMROZ Miryam, Institut Néel, CNRS (FRANCE)
ETTEMA Ad, Specs Nanotechnology (NETHERLANDS)
EYINK Kurt, Air Force Research Lab. (USA)
FAUVEL Véronique, Institut Néel, CNRS (FRANCE)
FEHRENBACH Bjoern, WEP (GERMANY)
FONTAINE Chantal, LAAS, CNRS (FRANCE)
FREY Alexander, Universitaet Wuerzburg, (GERMANY)
FRIGERI Paola, CNR-IMEM Institute (ITALY)
FURTMAYR Florian, Walter Schottky Institut (GERMANY)
FUSTER David, IMM, CSIC (SPAIN)
GACEVIC Zarko, ISOM (SPAIN)
GANZ Philipp, Karlsruhe Inst. of Technology (GERMANY)
GARCIA Basilio J., Univ. Autónoma de Madrid (SPAIN)
GARCÍA NUÑEZ Carlos, Univ. Autónoma de Madrid (SPAIN)
GASSLER Gerhard, Samtel Electron Devices (GERMANY)
GEMAIN Frédérique, CEA-Grenoble, LETI (FRANCE)
GENDRY Michel, Ecole Centrale de Lyon (FRANCE)
GENUIST Yann, Institut Néel, CNRS (FRANCE)
GÉRARD Lionel, Institut Néel, CNRS (FRANCE)
GHOLAMI MAYANI Maryam, Norwegian Univ. of Sc. & Tech. (NORWAY)
GILFERT Christian, University of Kassel (GERMANY)
GINÉS Mª LAIA, IMM, CSIC (SPAIN)
GIRAUD Etienne, EPFL (SWITZERLAND)
GLAS Frank, LPN, CNRS (FRANCE)

GOBAUT Benoit, Ecole Centrale de Lyon (FRANCE)
GOLLING Matthias, ETH Zürich (SWITZERLAND)
GONZALEZ Luisa, IMM, CSIC (SPAIN)
GONZÁLEZ Yolanda, IMM, CSIC (SPAIN)
GONZÁLEZ-POSADA Fernando, CEA-Grenoble, INAC (FRANCE)
GOUTARD Frédérick, RIBER (FRANCE)
GUINA Mircea, Tampere University (FINLAND)
GUZMÁN Alvaro, ISOM (SPAIN)
HAKKARAINEN Teemu, Tampere University (FINLAND)
HANANEL Ralph, AXT/ Geo Semiconductor (SWITZERLAND)
HATTON Carl, Veeco Instruments Inc. (USA)
HELFRICH Mathieu, Karlsruhe Inst. of Technol. (GERMANY)
HENNINGER Bernd, LayTec (GERMANY)
HERRANZ Jesús, IMM, CSIC (SPAIN)
HESTROFFER Karine, CEA-Grenoble, INAC (FRANCE)
HEYLENS Christophe, VBS Europe (BELGIUM)
HIERRO Adrian, ISOM (SPAIN)
HOFFMANN Dirk, TU Kaiserslautern (GERMANY)
HOPKINSON Mark, University of Sheffield (UK)
HUBER Frank, Huber Frank (GERMANY)
JO Masafumi, NIMS (JAPAN)
JOUANNEAU Romain, Vinci Technologies (FRANCE)
KAHOULI Abdelkarim, CRHEA, CNRS (FRANCE)
KAMPMEIER Jörn, University of Paderborn (GERMANY)
KASPRRZAK Jacek, Institut Néel, CNRS (FRANCE)
KAUFMANN Nils, EPFL (SWITZERLAND)
KERCKX Phillip, VBS Europe (BELGIUM)
KHAEMASAEVEE Patracha, Staib Instruments, Inc. (USA)
KHENG Kuntheak, CEA-Grenoble, INAC-UJF (FRANCE)
KLEMBT Sebastian, University of Bremen (GERMANY)
KLERKV Peter, DeMaCo Holland BV (NETHERLANDS)
KOTSAR Yulia, CEA-Grenoble, INAC (FRANCE)
KRAUS Andreas, TU Braunschweig (GERMANY)
KRUMRAIN Julian, Forschungszentrum Jülich (GERMANY)
KRUSE Carsten, University of Bremen (GERMANY)
LENZ Rudolf, EpiServe (GERMANY)
LEQUIEN Daniel, MCSE (FRANCE)
LETOUBLON Antoine, FOTON - INSA (FRANCE)
LI Ang, Istituto Nanoscienze-CNR (ITALY)
LISCHKA Klaus, Univ. Paderborn (GERMANY)
LOCHMANN Anatol, LayTec (GERMANY)
LOCKLEY John, AXT/ Geo Semiconductor (UK)
LUNA Esperanza, Paul-Drude-Institut (GERMANY)

LYAMKINA Anna, Rzhanov Institute (RUSSIA)
MALINVERNI Marco, EPFL (SWITZERLAND)
MARCADET Xavier, III-V Lab (FRANCE)
MARIETTE Henri, Institut Néel, CNRS (FRANCE)
MARTIN Denis, EPFL (SWITZERLAND)
MASSELINK W Ted, Humboldt University (GERMANY)
MAYER Bernd, Oclaro (SWITZERLAND)
MILLA Maria José, ISOM (SPAIN)
MONASTYRSKYI Grygorii, Humboldt University (GERMANY)
MONROY Eva, CEA-Grenoble, INAC (FRANCE)
MORIER-GENOUD François, EPFL (SWITZERLAND)
MUSSLER Gregor, Forschungszentrum Jülich (GERMANY)
NARITS(UK)A Shigeya, Meijo University (JAPAN)
NGUYEN THANH Tra, FOTON - INSA (FRANCE)
NOVAK Vit, Institute of Physics (CZECH REPUBLIC)
PENUELAS José, Ecole Centrale de Lyon (FRANCE)
PERLIN Piotr, UNIPRESS (POLAND)
PETROV Stanislav, SemiTEq JSC (RUSSIA)
PIERRET Aurélie, CEA-Grenoble, INAC (FRANCE)
POLIVKA Harald, Staib Instruments, Inc. (GERMANY)
PONCHET Anne, CEMES, CNRS (FRANCE)
PRESBERG Renaud, Vinci Technologies (FRANCE)
PRUDHOMMEAUX Elie, AZELIS electronics (FRANCE)
RABAROT Marc, INFICON (FRANCE)
REGRENY Philippe, Ecole Centrale de Lyon (FRANCE)
REICHL Christian, ETH Zürich (SWITZERLAND)
REUTER Alexandra, Karlsruhe Inst. of Technology (GERMANY)
RICHTER Mirja, IBM Zurich Research Lab (SWITZERLAND)
RIECHERT Henning, Paul-Drude-Institut (GERMANY)
RIEDL Hubert, Walter Schottky Institut (GERMANY)
RIEGER Torsten, Forschungszentrum Jülich (GERMANY)
RODRIGUEZ Jean-Baptiste, IES, CNRS (FRANCE)
RUDOLPH Andreas, Universität Regensburg (GERMANY)
RÜDT Christoph, Crea Tec Fischer & Co. (GERMANY)
RUIZ Ana, ICMM, CSIC (SPAIN)
SAINT-GIRONS Guillaume, Ecole Centrale de Lyon (FRANCE)
SAM-GIAO Diane, CEA-Grenoble, INAC (FRANCE)
SANGUINETTI Stefano, Università di Milano Bicocca (ITALY)
SCHMIDTBAUER Jan, Leibniz-Inst. for Crystal Growth (GERMANY)
SCHUBER Ralf, Karlsruhe Inst. of Technology (GERMANY)
SEIFRITZ Joerg, Omicron NanoTechnology (GERMANY)
SEMENOV Alexey, Ioffe Institute (RUSSIA)
SEMOND Fabrice, CRHEA, CNRS (FRANCE)

SIEKACZ Marcin, TopGaN Ltd. (POLAND)
SKIERBISZEWSKI Czeslaw, TopGaN Ltd. (POLAND)
SNAPI Noam, SCD (ISRAEL)
SOBANSKA Marta, Polish Academy of Science (POLAND)
SONG Yuxin, Chalmers Univ. of Technology (SWEDEN)
SONGMUANG Rudeesun, Institut Néel, CNRS (FRANCE)
SPRINGHOLZ Gunther, Johannes Kepler Univ. (AUSTRIA)
TARDE Cyril, RIBER (FRANCE)
TARNAWSKA Lidia, IHP (POLAND)
TATARENKO Serge, Institut Néel, CNRS (FRANCE)
TEISSEYRE Herryk, Polish Academy of Science (POLAND)
TIMOFEEV Vyacheslav, Rzhanov Institute (RUSSIA)
TOMITA Yuto, LayTec (GERMANY)
TONKIKH Alexander, Max Planck Institute (GERMANY)
TOURBOT Gabriel, CEA-Grenoble, INAC (FRANCE)
TOURNIÉ Eric, Univ. Montpellier 2 - CNRS (FRANCE)
TURSKI Henryk, UNIPRESS (POLAND)
UCCELLI Emanuele, EPFL (SWITZERLAND)
UGUR KATMIS Asli, Humboldt University (GERMANY)
ULLOA Jose Maria, ISOM (SPAIN)
UMANSKY Vladimir, Weizmann Institute of Science (ISRAEL)
UTZ Martin, Universität Regensburg (GERMANY)
UUSIPAIKKA Leena, DCA Instruments (FINLAND)
VANHATALO Jari, DCA Instruments (FINLAND)
VICO TRIVINO Noelia, EPFL (SWITZERLAND)
WALTHER Martin, Fraunhofer IAF (GERMANY)
WANG Deliang, Univ. of Sc. & Tech. of China (CHINA)
WANG Shumin, Chalmers Univ. of Technology (SWEDEN)
WASILEWSKI Zbig, NRC Canada (CANADA)
WEWIOR L(UK)asz, IMM, CSIC (SPAIN)
WIETLER Tobias F., Leibniz University (GERMANY)
WINKLER Achim, Veeco Instruments Inc. (GERMANY)
WOJNAR Piotr, Polish Academy of Science (POLAND)
ZHAO Huan, Chalmers Univ. of Technology (SWEDEN)
ZYTKIEWICZ Zbigniew, Polish Academy of Science (POLAND)

Veeco MBE:
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• Automated architecture
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• Built on ten years of reliable cluster tool technology
Learn more—
Stop by our
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Attend our
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Sunday evening, March 20
7:30 p.m., Hotel Le Pic Blanc
Visit
www.veeco.com/mbe_Euro2011

Registration Lobby Hotel Pic Blanc 11:00 – 20:00
Program of the 16 th European Molecular Beam Epitaxy Workshop
March 20 th -23 rd , 2011, Alpe d’Huez, France
Sunday, 20 th Monday, 21 st Tuesday, 22 nd Wednesday, 23 rd
18:00–19:15
Welcome
glass of
wine
Hotel Pic
Blanc
OPENING 8:15
Sanguinetti
(invited) 8:30
Mo1
Arsenides I
Guina
(invited) 8:30
Tu1
Arsenides
II
Ponchet Jo Wietler
(invited) 9:00
Richter
Golling
Ulloa Gilfert Saint-Girons
Hakkarainen Hopkinson Tonkikh
COFFEE BREAK
10:00 – 10:30
Novak
(invited) 10:30
Semenov
Luna
Mo2
Antimonides &
Phosphides
COFFEE BREAK
10:00 – 10:30
Rodriguez
(invited) 10:30
Bierwagen
(invited) 11:00
Tu2
New trends in
MBE
Borisova
Springholz
COFFEE BREAK
10:30 – 11:00
Eibelhuber
(invited) 11:00
Cornet Mussler Umansky
Desplanque Clarke Reichl
Walther Klembt Marcadet
Detz
BREAK 12:15 – 17:00
Lunch in a mountain
restaurant of the resort
+
Free time
Bonanni
(invited) 17:00
Furtmayr
(invited) 17:30
Petrov
Das
Cheze
Semond
Mo3
Nitrides
POSTER SESSION MoP1
19:00 – 20:45
19:30–21:30 Cocktail buffet:
VEECO
regional specialities
users’ in the exhibition rooms
meeting
Hotel Pic
Blanc 21:00
RIBER users' meeting
BREAK 12:15 – 17:00
Lunch in a mountain
restaurant of the resort
+
Free time
As
(invited) 17:00
Chauveau
(invited) 17:30
Tarnawska
Kotsar
Laumer
Siekacz
Tu3
Wide bandgap
POSTER SESSION TuP2
19:00 – 20:45
Cocktail buffet:
regional specialities
in the exhibition rooms
21:00
Workshop Banquet
Belahsene
We1
Group IV
Materials
We2
Devices
BREAK 12.45 -14.45
Lunch at Hotel Pic Blanc
Glas
(invited) 14:45
Uccelli
Li Ang
Dimakis
Alonso-Álvarez
We3
Nanowires
Hestroffer
Award ceremony
Closing address
16:30-17:00
Free time
Free access to keep-fit
center, swimming pool, spa
at Hotel Pic Blanc
Starting 19:00
Farewell Dinner