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TMS 2011, 140th Annual Meeting & Exhibition, Supplemental
Proceedings. Volume 2, Materials Fabrication, Properties,
Characterization and ModelingSupplemental Proceedings Volume
2:
TIMIS201 140th Annual Meeting & Exhibition
TIMIS201 140th Annual Meeting & Exhibition
Check out these new proceeding volumes from the TMS 2011 Annual
Meeting, available from publisher John Wiley & Sons:
2nd International Symposium on High-Temperature Metallurgical
Processing
EnergyTechnology 2011 : Carbon Dioxide and Other Greenhouse Gas
Reduction Metallurgy and Waste Heat Recovery
EPD Congress 2011
Light Metals 2011
Magnesium Technology 2011
Recycling of Electronic Waste II, Proceedings of the Second
Symposium
Sensors, Sampling and Simulation for Process Control
Shape Casting: Fourth International Symposium 2011
Supplemental Proceedings: Volume 1: Materials Processing and Energy
Materials
Supplemental Proceedings: Volume 2: Materials Fabrication,
Properties, Characterization, and Modeling
Supplemental Proceedings: Volume 3: General Paper Selections
To purchase any of these books, please visit www.wiley.com.
TMS members should visit www.tms.org to learn how to get discounts
on these or other books through Wiley.
Supplemental Proceedings Volume 2:
Materials Fabrication, Properties, Characterization, and
Modeling
About this volume The TMS 2011 Annual Meeting Supplemental
Proceedings, Volume 2: Materials Fabrication, Properties,
Characterization, and Modeling, is a collection of papers from the
2011 TMS Annual Meeting and Exhibition, held February 27-March 3,
in San Diego, California, U.S.A. The papers in this volume were
selected based on technical topic compatibility and represent
thirteen symposia from the meeting. This volume, along with the
other proceedings volumes published for the meeting, and archival
journals, such as Metallurgical and Materials Transactions and the
Journal of Electronic Materials, represents the available written
record of the 74 symposia held at the 2011 TMS Annual Meeting. The
individual papers presented within this proceedings volume have not
necessarily been edited or reviewed by the conference program
organizers and are presented "as is." The opinions and statements
expressed within the papers are those of the individual authors
only and are not necessarily those of anyone else associated with
the proceedings volume, the source conference, or TMS. No
confirmations or endorsements are intended or implied.
TIMIS201 140th Annual Meeting & Exhibition
®WILEY TIMS A John Wiley & Sons, Inc., Publication
,,i\t^^^^^^^^^m
Copyright © 2011 by The Minerals, Metals, & Materials Society.
All rights reserved.
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WILEY TIRAIS A John Wiley & Sons, Inc., Publication
Materials Fabrication, Properties, Characterization, and
Modeling
2011 Functional and Structural Nanomaterials: Fabrication,
Properties, Applications and
Implications
T. Nakamura, Y. Herbani, andS. Sato
Nanomaterials-Characteristics
Crystallization Kinetics and Giant Magneto Impedance Behavior of
FeCo Based Amorphous Wires 9
R. Roy, P. Sarkar, S. Singh, A. Panda, and A. Mitra
Sunday Evening Poster Session: Functional Materials
Fe-Based Amorphous-Nanocrystalline Thermal Spray Coatings 17 B.
Movahedi, and M. Enayati
Enhanced Photocatalytic Activity of Modified Ti02 for Degradation
of CH20 in Aqueous Suspension 25
H. Tonga, L. Zhaoc, D. Lia, andX. Zhanga
Preparation and Characterization of ZnS Thin Films Using Chemical
Bath Deposition Method: Effects of Deposition Time and Thermal
Treatment 43
W. Hsieh, K. Cheng, andS. Lue
Femtosecond Laser-Induced Synthesis of Colloidal AuAg Nanoalloys
from Aqueous Mixture of Metallic Ions 51
Y. Herbani, T. Nakamura, andS. Sato
v
Electrochemical Performances of Nanoporous Carbon Anode for Super
Lithium Ion Capacitor 59
Z Xiangyang, L. Shiju, Y. Juan, L. Changlin, and Z. Taikang
Effect of Temperature Schedule on the Particle Size of NiFe204
Spinel Nanopowder during Solid-State Reactions 67
Z. Zhang, G. Yao, Y. Liu, andJ. Du
Interfacial Properties of Cu-Nb Multilayers as a Function of
Dislocation/Disconnection Content 75
N. Abdolrahim, I. Mastorakos, H. Zbib, and D. Bahr
Long-Time Photoluminescence Kinetics in Quantum Dot Samples 83 K.
Krai, andM. Mensik
Synthesis and Characterization of Mullite 91 K. Paithankar, D.
Barbadikar, D. Peshwe, and A. Gandhi
Characterization of Hybrid Carbon-Nanotube Composite Interfaces as
a Function of Length Scale 99
H. Malecki, M. Duffy, S. Markkula, andM. Zupan
Synthesis and Characterization of Nanostrucrure Forsterite
Bioceramic for Tissue Engineering Applications 109
F. Tavangarian, R. Emadi, and M. Enayati
Investigation of Mechanical Properties of Silica/Epoxy
Nano-Composites by Molecular Dynamics and Finite Element Modeling
117
B. Mortazavi, J. Bar don, S. Ahzi, D. Ruch, and A. Laachachi
Tuesday Evening Poster Session: Ultra Fine Grained Materials
Basal-Plane Stacking-Fault Energies of Mg: A First-Principles Study
of Li- And Al-Alloying Effects 121
Z Jin, J. Han, X. Su, and Y. Zhu
Development of A1-TÍB2 Nanocomposite 129 Z Sadeghian, M. Enayati,
B. Lotfi, and P. Beiss
Dry Sliding Wear and Corrosion Behavior of Ultrafine-grained HSLA
Steel Processed using Multi Axial Forging 137
A. Padap, G. Chaudhari, andS. Nath
VI
Heterogenity of Microstructure Evolution in NiTi (50 at% Ni) Alloy
Severely Deformed by High Pressure Torsion 147
R. Singh, J. Fiebig, S. Ostendorp, H. Rösner, E. Prokofyev, R.
Valiev, S. Divinski, and G. Wilde
Aluminum Alloys: Fabrication, Characterization and
Applications
Development and Application
Hot Tensile Behaviour and Constitutive Analysis of
Al-5,5Zn-l,2Mg/Zr Alloys 157
P. Leo, E. Cerri, and H. McQueen
Production of Continuous Cast 3105 Coil-Stock for Thin Gauge Roller
Shutters 167
D. Spathis, andJ. Tsiros
Emerging Technologies
Preparation of Al-Li Alloys for Lithium-Air Secondary Battery by
Solid Diffusion Method 175
T. Cheng, Z. Lv, X. Zhai, M. Zhang, and G. Tu
Effects of Process Parameters on Rolled Precursor of Aluminum Foam
Sandwich Panel 179
B. Song, G. Yao, G. Zu, L. Wang, andZ. Guan
Preparation of Aluminum Foam Using a Novel Gas-Generating Agent 185
D. Huo, X. Zhou, T. Zhang, J. Qin, J. Li, and H. Zhao
High Temperature Dry Sliding Wear Behaviour of Aluminium-Silicon /
Graphite Composite Processed by Stir Casting 191
G. Rajar am, S. Kumar an, T. Rao, andM. Kamaraj
Preparation and Characterization of Short Carbon Fiber Reinforced
Aluminium Matrix Composites 199
P. Yan, G. Yao, J. Shi, X. Sun, andG. Lv
Vll
Materials Characterization Effect of Ultrasonic Impact Treatment on
a 5456 Aluminum Alloy Characterized through Micro-Specimen Testing
and X-Ray Tomography 205
C. Scheck, K. Tran, C. Cheng, and M. Zupan
Failure Loads and Deformation in 6061-T6 Aluminum Alloy Spot Welds
213 R. Florea, K. Solanki, D. Bammann, B. Jordon, and M.
Castanier
Numerical Modeling
Modeling Performance of Protection Materials Aluminum 7020-T651 and
Steel 221
J. Chinella
Comprehensive Thermo-Mechanical Validation of Extrusion Simulation
Cycle for Al 1100 Using HyperXtrude 229
A. Parkar, C. Bouvard, S. Horstemeyer, E. Marin, P. Wang, and M.
Horstemeyer
Mechanical Properties and Casting Characteristics of the Secondary
Aluminum Alloy AlSi9Cu3(Fe) (A226) 237
P. Pucher, H. Böttcher, H. Kaufmann, H. Antrekowitsch, and P.
Uggowitzer
Comparison of Different FEM Codes Approach for Extrusion Process
Analysis 245
L. Donati, L. Tomesani, N. Khalifa, and A. Tekkaya
Numerical Prediction of Grain Shape Evolution during Extrusion of
AA6082 Alloy 253
A. Segatori, L. Donati, andL. Tomesani
Analysis of Charge Weld Evolution for a Multi-Hole Extrusion Die
263 A. Segatori, L. Donati, B. Reggiani, andL. Tomesani
Solidification
Vlll
Solidification Analysis of Al-Si Alloys Modified with Addition of
Cu Using In- Situ Neutron Diffraction 279
D. Sediako, W. Kasprzak, I. Swainson, and O. Garlea
Novel Grain Refiner for Al-Si Alloys 291 M. Nowak, andN. Babu
Application of Neutron Diffraction in Analysis of Residual Stress
Profile in the Cylinder Web Region of as-Cast V6 Aluminum Engine
Block with Cast-In Iron Liners 299
D. Sediako, R. Ravindran, C. Hubbard, F. D'Elia, A. Lombardi, A.
Machin, and R. Mackay
Effects of A1-8B Grain Refiner on the Structure, Hardness and
Tensile Properties of a New Developed Super High-Strength Aluminum
Alloy 309
M. Alipour, M. Emamy, J. Rasizadeh, M. Karamouz, and M.
Azarbarmas
Thermal Mechanical Processing
Study of the Artificial Aging Kinetics of Different AA6013-T4 Heat
Treatment Conditions 321
J. Berneder, R. Prillhofer, J. Enser, P. Schulz, and C.
Melzer
Estimating Response to Hot Rolling of Al-Mn-Mg Alloys from Hot
Torsion Testing 329
H. McQueen
P. Leo, E. Cerri, andH. McQueen
IX
Characterization and Processing Techniques for Composites
Thermo-Mechanical Behavior of Hdpe/Sugarcane Bagasse
Fiber/Organoclay Nanocomposites 349
A. Castillo, A. Teran, A. Chinellato, M. Nascimento, F. Diaz, and
E. Moura
Development of New Composite Materials
Machinable Aluminum Matrix Composite 359 W. Harrigan
Stability and Lithium Adsorption Property of LiMn204-LiSb03
Composite in Aqueous Medium 367
X. Shi, L. Ma, B. Chen, H. Xu, X. Yang, and K. Zhang
Reinforced Steel/Polymer/Steel Sandwich Composites with Improved
Properties 375
H. Palkowski, O. Sokolova, and A. Carrada
Understanding Composite Performance
K. Yanase, andJ. Ju
Modelling Shear Fracture of Hybrid CFRP/Ti Laminates with Cohesive
Elements; Effects of Geometry and Material Properties 391
P. Naghipour, M. Bartsch, J. Hausmann, andK. Schulze
x
Computational Thermodynamics and Kinetics
Brent Fultz Honorary Session II Phonon Thermodynamics of Binary Fe
Alloys 401
M. Lucas
Defects: Thermodynamics and Kinetics of Grain Boundaries,
Interfaces, Surfaces and Dislocations
Phase-Field Simulation of Segregation to Stacking Fault (Suzuki
Effect) in Co- Ni Based Superalloy 409
Y. Koizumi, S. Suzuki, T. Otomo, S. Kurosu, Y. Li, H. Matsumoto,
and A. Chiba
Microstructual Evolution
Phase-Field Simulations of Bainitic Phase Transformation in 100Cr6
417 W. Song, U. Prahl, W. Bleck, and K. Mukherjee
Microstructure Evolution and Analysis of Single Crystal
Nickel-Based Superalloy during Compression Creep 427
Z Shu, T. Sugui, L. Fushui, L. Anan, andL. Jingjing
Poster Session: Computational Thermodynamics and Kinetics of
Materials
The Application of Thermodynamic Analysis in Preparing the MnZn
Ferrites Precursor 435
X. Ping, Y. Yaohua, Z. Peiyu, and C. Xiaofang
Phase Equilibria of the La-Ni-Cu Ternary System at 673 K:
Thermodynamic Modeling and Experimental Validation 441
X. An, Q. Li, J. Zhang, S. Chen, and Y. Yang
Statistical Model of Precipitation Kinetics for Recycled Commercial
Aluminum Alloys 449
Z Liu, V. Mohles, O. Engler, and G. Gottstein
XI
Thermodynamics Calculation of CuO-NH3+NH4Cl Solution System 457 W.
Zheng, D. Li, Z. Xiao, Q. Chen, and H. Tong
Development of Accurate Models for the Microstructure and
Properties of Molten Salts 461
A. Gray-Weale, P. Masset, and A. Jacob
A Heat Management Model for Hardness Uniformity of Multi-Pass Laser
Heat Treatment Using Direct Diode Laser 469
S. Santhanakrishnan, and R. Kovacevic
Thermodynamics, Phase Stability and Phase Transformations
Thermomechanical Processing Design of Nanoprecipitate Strengthened
Alloys Employing Genetic Algorithms 477
P. Rivera-Diaz-del-Castillo, Maarten de Jong, and M. Sluiter
David Pope Honorary Symposium on Fundamentals of Deformation and
Fracture of Advanced Metallic
Materials
Deformation, Fracture, and Advanced Characterization
Techniques
Intelligent Microscopy for the Study of Fracture and Fatigue 489 D.
Fullwood, B. Adams, T. Rampton, and A. Khosravani
Deformation, Fracture, and Hydrogen Effects
Influence of Hydrogen Loading on the Tensile Behavior of Fe-Ga
Alloys 497 M. Ramanathan, B. Saha, C. Ren, G. Garside, andS.
Guruswamy
XI1
B. Biner, and L. Kubin
Geometrical Construction and Structure of Quasi-Periodic Grain
Boundaries in Cubic Materials 513
M. Shamsuzzoha
Influences of Material and Process Parameters on Delayed Fracture
in TR1P- Aided Austenitic Stainless Steels 521
X. Guo, and W. Bleck
Intermetallics I
T. Takasugi, and Y. Kaneno
Intermetallics II and Ti alloys
Some Unusual Aspects of the Deformation Of FeAl and Fe2MnAl 537 /.
Baker
Recent Progress in High Temperature TiAl Alloys 547 G. Chen, L.
Zhao, J. Lin, andX. Xu
Intermetallics III, Superalloys, and Gum Metal
Overview of Creep Deformation of Nickel Base Superalloys and
Intermetallics 557
D. Shah
Localized Shear Deformation in Gum Metal at Ideal Strength 567 S.
Kuramoto, T. Furuta, N. Nagasako, andJ. Morris
xiu
Session I
The Effect of Crystallographic Orientation on Void Growth: A
Molecular Dynamics Study 577
M Bhatia, K. Solanki, A. Moitra, and M. Tschopp
Room Temperature Creep and Substructure Formation in Pure Aluminum
at Ultra-Low Strain Rates 585
& Junjie, I. Ken-ichi, H. Satoshi, andN. Hideharu
Session II
Development of <111> Fiber Texture and {111 }<112>
Shear Bands in Pure Al Metal by Wire Drawing 593
M. Shamsuzzoha
N. Konchakova, R. Mueller, F. Barth, F. Balle, andD. Eifler
Role of Austenite Plasticity in the Deformation of Superelastic
Nitinol 609 D.Xu, andR. Ritchie
Vanadium Effects on a BCC Iron Sigma 3 (111) [1-10] Grain Boundary
Strength 617
S. Kim, S. Kim, and M. Horstemeyer
Fracture Behavior of Short Carbon Fiber Reinforced Aluminium Matrix
Composite 621
P. Yan, G. Yao, J. Shi, andX. Sun
Session III
Stress Intensity Factor Solutions for Friction Stir Spot Welds of
Magnesium AZ31 Alloy 627
T. Tang, M. Horstemeyer, B. Jordan, and P. Wang
xiv
Deformation Induced Phase Transformation during Machining of
Ti-5553 633 D. Y an, G. Littlefair, and T. Pasang
Fatigue and Corrosion Damage in Metallic Materials: Fundamentals,
Modeling and Prevention
Fatigue and Corrosion Interaction and Materials Corrosion
Effect of Proximity and Dimension of Two Artificial Pitting Holes
on the Fatigue Endurance of Aluminum Alloy 6061-T6 under Rotating
Bending Fatigue Tests 643
G Almaraz, V. Lemus, andJ. López
Fatigue of Nanocrystalline Materials and Fatigue Property
Enhancement
Research on HCF Tests and Damage Model of TCI 1 Alloy Welded Joints
....651 X. Liu, and G. H ai-ding
Fatigue Behavior of Al 6082-T4 and Al 7075-T73 after Ball
Burnishing 659 Y. Fouad, M. Mhaede, andL. Wagner
Fatigue Propertv-Microstructure Relationships and Crack
Growth
A Modified LEFM Approach for the Prediction of the Notch Effect in
Fatigue 667
M. Endo, K. Yanase, S. Ikeda, and A. McEvily
Resistivity Based Evaluation of the Fatigue Behaviour of Cast Iron
675 H. Germann, P. Starke, and D. Eifler
Microstructure-Sensitive Probabilistic Fatigue Modeling of Notched
Components 683
W. Musinski, andD. McDowell
M. Sadawy
Effect of Temperature on the Loss of Ductility of S-135 Grade Drill
Pipe Steel and Characterization of Corrosion Products in C02
Containing Environment 699
A. Bajvani Gavanluei, B. Mishra, and D. Olson
Corrosion Behavior and Galvanic Corrosion Studies of TÍ-6A1-4V
Alloy GTA Weldment in HC1 Solution 707
M. Atapour, E. Mohammadi Zahrani, M. Shamanian, and M. Fathi
Comparative Study of Hot Corrosion Behavior of Plasma Sprayed
Yttria and Ceria Stabilized Zirconia Thermal Barrier Coatings in
Na2S04+V205 at 1050°C 715
M Mahdipoor, M. Rahimipour, and M Habibi
The Effect of Temperature on the Corrosion Behavior of 625
Superalloy in PbS04-Pb305-PbCl-ZnO Molten Salt System with 10 wt. %
CdO 725
E. Mohammadi Zahrani, and A. Alfantazi
Frontiers in Solidification Science
Posters
A. Meysami, R. Ghasemzadeh, H. Seyedyn, M. Aboutalebi, andR.
Rezaei
xvi
A Numerical Benchmark on the Prediction of Macrosegregation in
Binary Alloys 755
H. Combeau, M. Bellet, Y. Fautrelle, D. Gobin, E. Arquis, O.
Budenkova, B. Dussoubs, Y. Duterrail, A. Kumar, B. Goyeau, S.
Mosbah, T. Quatravaux, M. Rady, C. Gandin, and M. Zaloznik
ICME: Overcoming Barriers and Streamlining the Transition of
Advanced Technologies to Engineering
Practice - The 12th MPMD Global Innovations Symposium
Emerging and Fundamental Techniques and the Advancement of ICME in
Industry
Modeling and Simulation of Mechanical Properties of Magnesium Alloy
Wheel Casting for Automobile 765
L. Huo, Z. Han, X. Zhu, J. Duan, A. Wang, andB. Liu
Modeling and Simulation Tools
K. Ferris, and D. Jones
Massively Parallel Simulations of Materials Response
Session II
Lights - Open Source Discrete Element Simulations of Granular
Materials Based on Lamps 781
C. Kloss, and C. Gonina
XVll
Session III Atomic Scale Deformation Mechanisms of Amorphous
Polyethylene under Tensile Loading 789
M. Tschopp, J. Bouvard, D. Ward, andM. Horstemeyer
Recent Developments in the Processing, Characterization, Properties
and Performance of
Metal Matrix Composites
General and Nano-Composites
Low Density Magnesium Matrix Syntactic Foams 797 J. DeFouw, and P.
Rohatgi
Joining of Advanced Aluminum-Graphite Composite 805 N. Hung, M.
Velamati, M. Garza-Castañon, E. Aguilar, and M. Powers
Multimodal, Processing and Microstructure
Effect of A1+B4C Agglomerate Size on Mechanical Properties of
Trimodal Aluminum Metal Matrix Composites 813
B. Yao, T. Patterson, Y. Sohn, M. Shaeffer, C. Smith, M. van den
Bergh, and K. Cho
Effects of S PS Parameters on the Mechanical Properties and
Microstructures of Titanium Reinforced with Multi-Wall Carbon
Nanotubes Produced by Hot Extrusion 821
T. Threrujirapapong, K. Kondoh, J. Umeda, B. Fugetsu, and T.
Mimoto
xvni
Microstructural Development of Al-15wt.%Mg2Si In Situ Composite
with Be Addition 829
M Azarbarmas, M. Emamy, J. Rasizadeh, M. Alipour, and M.
Karamouz
Microstructural Properties and Wear Behaviour of AlSi9Mg Matrix
B4CP Reinforced Composites 837
F. Top tan, I. Kerti, A. Sagin, M. Cigdem, S. Daglilar, and F.
Yuksel
Modification of Al-Mg2Si In Situ Composite by Boron 843 M.
Azarbarmas, M. Emamy, J. Rasizadeh, M. Karamouz, and M.
Alipour
In-Situ Synthesis of A1N/Mg Matrix Composites 851 X. Ma, S. Kuplin,
D. Johnson, andK. Trumble
Performance Evaluation of Particulate Reinforced Al-SiC Bolted
Joints 859 G. William, S. Shoukry, andJ. Prucz
Processing, Microstructure and Mechanical Properties II
Effect of MgAl204 on the Superficial Hardness of Hybrid-Multimodal
Al/SiC Composites Processed by Reactive Infiltration 867
M. Montoya-Davila, M. Pech-Canul, andR. Escalera-Lozano
Corrosion and Wear Behaviour of Aluminum Alloy 6061-Fly Ash
Composites 873
A. Bhandakkar, B. Balaji, R. P ras ad, andS. Sas try
Interface Evolution in Tungsten Wire Reinforced Stainless Steel
Composites 883
P. Kumar, andM. Krai
Effects of Annealing on the Growth Behavior of Intermetallic
Compounds on the Interface of Copper/Aluminum Clad Metal Sheets
895
L. Xiaobing, Z. Guoyin, andD. Qiang
xix
Surfaces and Heterostructures at Nano- or Micro- Scale and Their
Characterization, Properties, and
Applications
Coatings, Surfaces, and Interfaces II - and - Magnetic
Heterostructures I
Application of the Strong Contrast Technique to Thermoelastic
Characterization ofNanocomposites 905
M. Baniassadi, A. Ghazavizadeh, D. Ruch, Y. Rémond, S. Ahzi, and H.
Garmestani
Energy and Catalysis Technologies II - and - Biological
Applications
Colloid-Chemical Nanoprocesses and Nanotechnologies on the Basis of
Oxyhydrate Systems of Rare-Earth Elements 911
T. Prolubnikova, Y. Sucharev, T. Ukolkina, and K. Nosov
Thermally Activated Processes in Plastic Deformation
Deformation Mechanisms and Polvcrystal Plasticity
Comparative Hot-Work Constitutive Analyses Of Carbon/HSLA and
Stainless Steels with Linkage to Microstructural Evolution
921
H. McQueen, Y. Li, I. Rieiro, M. Carsi, and O. Ruano
Grain Boundary Evolution and Dislocation Core Effects
Experimental Measurements of the Shear-Coupled Stress Driven Grain
Boundary Migration in Al Bicrystals 931
D. Molodov, T. Gorkaya, andG. Gottstein
xx
Smelting and Reduction Processes
Experimental Study on Reduction Roasting and Separation of
Nickeliferous Latente by Microwave Heating 941
L. Yi, Z. Huang, B. Hu, X. Wang, and T. Jiang
Author Index 953
Subject Index 959
-i-
Fabrication, Properties, Applications and
Implications
The proceedings contained in this section have not been edited or
reviewed by the conference program organizers. The opinions and
statements expressed in the proceedings are those of the authors
only and are not necessarily those of the editors or TMS staff. No
confirmations or endorsements are intended or implied.
This page intentionally left blank
Supplemental Proceedings: Volume 2: Materials Fabrication,
Properties, Characterization, and Modeling TMS (The Minerals,
Metals & Materials Society), 2011
FABRICATION OF GOLD-PLATINUM NANO ALLOY BY HIGH- INTENSITY LASER
IRRADIATION OF SOLUTION
Takahiro Nakamura, Yuliati Herbani, Shunichi Sato
Institute of Multidisciplinary Research for Advanced Materials,
Tohoku University Katahira 2-1-1, Aoba-ku, Sendai 980-8577,
Japan
Keywords: Femtosecond laser, Liquid, Au-Pt nanoalloy
Abstract
Gold-platinum (Au-Pt) solid solution nanoalloys were fabricated by
high-intensity femtosecond laser irradiation of mixed solution of
auric and platinum ions. Photo-absorption spectra of prepared
solutions were measured by UV-visible spectrophotometer before and
after irradiation. The fabricated particles were characterized by
TEM and XRD. While two representative diffraction peaks are
commonly observed between peak the positions of pure bulk gold and
platinum for bulk because of a large immiscibility gap in a Au-Pt
binary system, only a single diffraction peak was detected for
single-nanometer sized Au-Pt nanoalloy particles fabricated by
high-intensity laser irradiation of mixed solution of auric and
platinum ions with the concentration of 5.0χ10"4 Μ. This finding
demonstrates that solid solution Au-Pt nanoalloys are successfully
fabricated only by high-intensity laser irradiation of aqueous
solution without any chemicals.
Introduction
Binary alloy nanoparticles (NPs) have been intensively studied
especially in the research field of catalysis1'2 because of their
bifunctional catalytic properties. Currently, gold-platinum (Au-Pt)
nanoalloys are attracted much attention for electrocatalysis in a
fuel cell3'4. The Au-Pt nanoalloys are expected to provide
synergistic catalytic activities such as suppression of adsorbed
poisonous species like carbon monoxide (CO) on Pt atoms, and the
change in electronic band structure to modify the strength of the
surface adsorption. The decreases of activation energy promoting
oxidative desorption and suppressing the adsorption of CO was
considered as a factor that leads to a sufficiently high
adsorptivity to support catalytic oxidation in alkaline
electrolytes5'6. Au-Pt nanoalloys are prepared mainly by chemical
processes7"13 in a form supported on a specially prepared substrate
such as SÍO2" and carbon 2' 13 to date. The process commonly needs
a series of complex procedures and often uses some chemicals that
might be highly reactive and cause environmental and biological
problems.
Recently, we have demonstrated a method for the preparation of
metal NPs of gold14, platinum15 and silver by using high intensity
laser irradiation of the metal ion solution. This technique is
expected to produce many kinds of metal and their alloy NPs
directly in the solution without any complex procedures and harmful
chemicals. In this study, we describe the fabrication of Au-Pt
nanoalloy in a mixed solution of auric and platinum ions by high
intensity laser irradiation of the solution. Effects of the
fraction of auric and platinum ions in the solution on the
composition and structure of Au-Pt nanoalloys were investigated.
The fabrication mechanism of the NPs was also discussed.
Experimental
3
Mixed solutions of auric and platinum ions with different fraction
were prepared by the following procedure. Auric and platinum
aqueous solutions were separately prepared by dissolving hydrogen
tetrachloroauric (III) tetrahydrate powder (HAUCI33H2O, Wako Pure
Chemical Industries, Ltd., > 99.9 %) and hydrogen
hexachloroplatinic (IV) hexahydrate powder (Η2Ρθ6·6Η20,
Sigma-Aldrich Co., > 99.9 %) in extra-pure water. The
concentration of each solution was set to 5.0X10"4 M. Subsequently,
both solutions were mixed with different molar fractions. Samples
are labeled by the molar fraction of auric and platinum ions. For
example, 50 % of auric and platinum solution is labeled as
Au50Pt50. All the solutions were transparent, and no apparent
difference was observed. Figure 1(a) shows UV-visible absorption
spectra of prepared solutions with different molar fraction of
auric and platinum ions measured by a UV- visible spectrophotometer
(JASCO Co., V630 iRM). UV-visible absorption spectrum was shifted
from that of auric (AulOOPtO) to platinum solution (AuOPtlOO) with
decreasing the fraction of auric ion in the solution. As a target
of laser irradiation, 3 milliliters of each aqueous solution was
dispensed in a 10x10x45 mm quartz glass cuvette that is optically
transparent at the wavelength of incident laser light. Femtosecond
laser beam was generated from a chirped-pulse amplified Ti:sapphire
laser system with the wavelength of 800 nm. The pulse energy was 5
mJ with the pulse width of 100 fs and the repetition rate was 30
Hz. The laser beam was introduced to the cuvette normal to its
surface and tightly focused in the solution by an aspheric lens
with the focal length of 8 mm and the numerical aperture of 0.5.
The spot diameter was estimated to be 175 μηι in a diameter.
Theoretical estimation of the laser intensity was 2.1*1014W/cm2
taking into account that a laser beam radius is 3.2 mm before the
focusing lens, and the refractive index of the solution is 1.33
(water). The irradiation time was set to 30 min in every
experiment. Optical characteristics of the solution after laser
irradiation were evaluated by a UV-visible spectrometer.
Transmission electron microscopes (TEM: JEOL, JEM2000EXII) were
employed to take electron micrographs of the products after
irradiation. The samples for TEM observation were prepared by
falling a few drops of the solution on a carbon-coated copper grid
(Okenshoji Co., Ltd., Micro grid type-B) immediately after the
irradiation and dried in air at room temperature. The samples for
the XRD measurement were prepared by freeze-drying and placing the
obtained powder on a non-reflecting single crystal silicon plate
(Rigakti Co.), which is specially made to avoid any diffraction
peak of silicon over measurement range.
Figure 1. Uv-vis. absorption spectra of the mixed solution of auric
and platinum ions with different fractionsf (a) before and (b)
after irradiation.
4
Results
A tiny flash of luminescence and fine bubbles were observed around
the focal point during laser irradiation. These gasses were
identified as oxygen and hydrogen by Chromatographie analysis
(GC-8A, Shimadzu Co.). The gases were probably produced by the
decomposition of water molecules through the laser induced break
down16 facilitated by a high intensity laser field. The
transparency of the solution gradually changed during the laser
irradiation and resultant color of the solution after 30 minutes
irradiation strongly depended on the molar fraction of auric and
platinum ions in the solution; red-purple for AulOOPtO and light-
brown for AuOPt 100.
Figure 1(b) shows a representative set of UV-visible absorption
spectra of the solutions with different molar fractions after
irradiation. The spectra were measured promptly after the
irradiation. In the spectrum of auric solution (AulOOPtO), an
absorption peak at 520 nm was observed arising from surface plasmon
resonance (SPR) of gold nanoparticles. The peak position shifted to
shorter wavelength, and the absorbance decreased with the decrease
in the fraction of auric ion in the solution.
TEM bright field images of the particles are shown in Fig. 2. Mean
particle size of each sample evaluated from the TEM images was also
shown below the micrograph. As seen in the figure, particle size in
the micrographs became smaller with the decrease in the fraction of
auric ion in the solutions. This result is comparable to the fact
that gold particles tend to grow and crystallize faster than other
noble metals such as palladium and platinum because of its property
of low melting point (1336 K) and no affinity to oxygen.
Figure 2. TEM images of the NPs fabricated by high intensity laser
irradiation of mixed solution of auric and platinum ions with
different fractions.
5
To determine the structural characteristics of the fabricated
particles, XRD measurement were employed for all samples. A
representative set of profiles is shown in Fig. 3. The typical XRD
peak positions of gold and platinum from 1 1 1 planes are also
indicated by broken lines for comparison. As seen in the figure,
the diffraction patterns of the particles in AulOOPtO and AuOPtlOO
are indexed to be an fcc-type cubic lattice of bulk gold and
platinum. XRD peaks in the profile were shifted from the peak
position of gold to that of platinum with decreasing the fraction
of auric ion in the solution. The results from the structural
analysis of the fabricated particles by using Integrated X-ray
Powder Diffraction Software (Rigaku Co.) are summarized in Table 1.
Crystalline sizes of the particles calculated by Scherrer's
equation seemed to be larger than the particle sizes observed in
TEM images (Fig. 2). This might be arising from crystal growth
during sample preparation by freeze-drying. The crystalline sizes
varied from 50 nm to 6 nm with the decrease in the composition of
auric ion in the solution. This result denotes the same tendency as
the result from TEM observation (Fig. 2). Lattice constants of the
particles fabricated in the solutions of AulOOPtO (a = 4.082 Â) and
AuOPtlOO (a = 3.927 Â) were in a good agreement with those of bulk
gold and platinum. Interestingly, lattice constant of the
fabricated nanoalloy was almost linearly changed from that of bulk
gold to platinum depending on the fraction of auric and platinum
ions in the mixed solutions. This result clearly indicates the
solid solution Au-Pt nanoalloys with intended composition were
successfully fabricated in the solutions only by high-intensity
laser irradiation of solutions without any chemical.
Figure 3. XRD profiles of the NPs fabricated by high intensity
laser irradiation of mixed solution of auric and platinum ions with
different fraction.
Table 1. Characteristic parameters of nanoparticles evaluated from
XRD peaks Solution
AulOOPtO
Au90Pt10
Au80Pt20
Au70Pt30
Au60Pt40
Au50Pt50
AU40R60
Au30Pt70
Au20Pt80
Au10Pt90
AuOPtlOO
504.9
149.2
88.5
69.8
59.6
61.6
59.4
66.8
84.7
106.3
144.1
6
Discussions
The mechanism of the formation of Au-Pt nanoalloys by laser
irradiation of the solution without using any reducing agent was
attributed to the optically induced decomposition of water
molecule. Namely, solvated electrons and hydrogen radicals were
formed in the aqueous solution like a kind of photochemical
reaction'7"19. Generation of oxygen and hydrogen gasses around the
focal spot during laser irradiation16 strongly indicates that
hydrogen and hydroxyl radicals were simultaneously produced in the
solution. Among them, solvated electrons and hydrogen radicals can
act as a strong reducing agent in the solution. Therefore, metal
ions in the solution were easily reduced to zero-valance atoms
forming 'core' of the particles. When the size of the particles
reached several nanometers, most of the atoms produced by the laser
irradiation had been expended and the growth of the particles was
ceased. The binary phase Au-Pt alloy was generally produced in bulk
because of an immiscible gap for Au-Pt bulk alloy appeared in the
phase diagram. G. C. Bond pointed out that the gap arises from the
difference in the electronic energy levels of gold and platinum,
and small particles highly tend to form homogeneous alloys because
all the atoms retain their electronic structure, and hence no
rehybridization due to band formation takes place20. It is also
reported elsewhere7"13 that Au-Pt alloy nanoparticles with a
diameter of single nanometer are chemically synthesized in all
composition range. In fact, the particle size of nanoalloy measured
in our study was single nanometer and the compositions of resultant
Au-Pt nanoalloys showed a relatively good agreement with the molar
fraction of solution. This is caused by strong reducing power of
solvated electrons and hydrogen radicals produced by high-intensity
laser irradiation of aqueous solution. In fact, solid solution
Au-Pt nanoalloys are successfully fabricated only by high-intensity
laser irradiation of aqueous solution without any chemicals.
Conclusion
We have demonstrated the fabrication of solid solution Au-Pt
nanoalloy with regulated compositions in a high intensity optical
field produced by tightly focused femtosecond laser pulses in a
mixed solution of auric and platinum ions. It should be noticed
that the technique is quite simple and 'green' process without
using any chemicals except for metal salt. In addition, it is
applicable to other kinds of binary and ternary system.
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7
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8
CRYSTALLIZATION KINETICS AND GIANT MAGNETO IMPEDANCE BEHAVIOR OF
FeCo BASED AMORPHOUS WIRES
R.K. Roy, P. Sarkar, S. Singh, A. K. Panda, A. Mitra
National Metallurgical Laboratory (CSIR); Burma Mines; Jamshedpur,
Jharkhand 831007, INDIA.
Keywords: Amorphous wire, Crystallization Kinetics, Thermal
Stability, GMI Behavior
Abstract
The effects of Nb addition on crystallization kinetics and giant
magneto impedance (GMI) properties of F^Q^çSisBn amorphous wires
prepared by in-water quenching system have been investigated.
Thermal behaviors of the wires have been investigated by thermal
electrical resistivity measurement and differential scanning
calroimetry. The substitution of 4 at% Nb for Fe and Co increases
crystallization temperature and merges two crystallization peaks
into one peak, leading to a significant increase in thermal
stability against crystallization for Fe37Co37Nb4SigBi4 wire. The
formation of Fe2Nb phase due to addition of Nb increases the
activation energy for crystallization from 425 to 550 kl/mol. The
GMI properties of the alloys are evaluated at driving current
amplitude of 10 mA and a frequency of 400 kHz. The alloys show the
single peak behavior in the GMI profile. The change in GMI
properties increases from 10% at 0 at% Nb to 25% at 4 at% Nb.
Introduction
Since the discovery of the GMI effect in Co-based amorphous wires
in 1994, the ferromagnetic amorphous wires are widely used in
various magnetic sensors such as antitheft systems, magnetic
marking and labeling, geomagnetic measurements, space research,
target detection and tracking [1-3]. Due to high demand in the
field engineering and industrial sectors, a large number of
research works have been carried out for the improvement of GMI
sensors. The main focus is on the development of new materials and
subsequent processing of the materials by thermal treatment and/or
tensile loading. Therefore, a thorough understanding of GMI
phenomena with respect to alloy compositions and annealing
temperature, time, tensile stress have a great emphasis for
developing novel magnetic sensors.
The water-quenched amorphous wire preparation is dependent on three
factors, i.e., (i) solidification of the metallic melt stream at
high cooling rates and within the stable distance from the ejection
point, (ii) use of a cooling fluid with low viscosity and surface
tension, and (iii) stable and non-turbulent flow of the cooling
liquid at high velocities [4]. Amongst these factors, first and
second points are processing parameters and third point is
dependent on alloy composition [5], Therefore, alloy compositions
should be optimized for stable and non-turbulent flow of the alloys
in water, resulting in the production of defect-free and continuous
wire. The optimized alloy compositions are also responsible for
high GMI effect in the wires [4].
Since amorphous alloys are thermodynamically instable, the physical
properties of the alloys are frequently changed with respect to
both temperature and time. It hindrances amorphous alloys
9
used in practical applications. However, the alloys become
thermodynamically equilibrium after structural relaxation and
nanocrystallization at higher temperature [6]. The crystallization
kinetics not only changes thermal behavior of amorphous alloys but
also influences magnetic and GMI properties at different conditions
[7, 8]. Moreover, the controlled crystallization causes the
tailoring of the microstructure, resulting in the desired
properties in nanocrystalline-amorphous matrix alloys [9, 10].
Therefore, the studies on thermal stability and crystallization
kinetics of the amorphous alloys are important for its practical
application. Despite several studies published in the literature
about FeCo-based amorphous wires and their GMI properties, the
crystallization behavior study of the wires is very little. The aim
of this work is to present the effect of Nb on crystallization
kinetics of the FeCoBSi based amorphous wire and subsequently the
effect of structural changes on GMI properties.
Experimental Procedure
The ingots with a nominal composition of Fe39Co39SisBi4 (FC1) and
Fe37Co37Nb4SigBi4 (FC4) were prepared by arc melting the mixture of
pure elements (>99.9 wt %) in an argon atmosphere. The amorphous
wires of the alloys were produced by in-water quenching technique
(Figure la). In this process, the small pieces of ingot were
remelted in a quartz crucible with a nozzle diameter of 150 μπι and
ejected on the water of rotating drum through the nozzle under
argon gas pressure of 300 kPa (Figure lb). The amorphous wires
produced from this method are 120 μπι in diameter and 3-4 m in
length. The structures of as-cast and annealed wires were
characterized by x-ray diffractometer (XRD) using CuKa radiation
(λ= 0.1540 nm). The crystallization kinetics was investigated at
different heating rates of 20, 30, 40 and 50°C/min by differential
scanning calorimetry (DSC) with a Perkin-Elmer Diamond DSC under a
continuous flow of purified argon. Electrical resisitivity
measurement was done using thermal electrical resistivity (TER)
unit of Ulvac-Riko with a heating rate of 10°C/min. The
magneto-impedance was measured by the four probe technique where
the driving field was generated by passing an ac current and the
system was capable of generating current amplitude (Iac) ranging
between 1-20 mA with a maximum frequency 2000 kHz. A Helmholtz coil
was used to apply a dc external magnetic field parallel to the axis
of the sample. The percentage of GMI ratio (ΔΖ/Ζ) has been
calculated from the first harmonic signal using the relation
ΔΖ fZ(H) - Z(H0}1 —% = I , , x loo
where HQ= 0 kAm-1, the minimum dc applied field.
10
Figure 1. (a) Wire preparation by in-water quenching technique, (b)
Schematic diagram of cross- sectional view of the technique.
Results and Discussion
3.1 Structure of Water Quenched Wires
Figure 2 shows the surface smoothness and structure of as-cast
wire. The water-quenched as- cast represents quite smooth surface,
as observed by scanning electron microscopy (SEM) image (Figure
2a). The smooth surface of the wire is dependent on the
optimization of process parameters and alloy compositions. The
structure of as-cast wires is basically amorphous in nature,
observing by a halo diffraction peak and no appreciable crystalline
peaks (Figure 2b).
Figure 2. As-cast wires showing (a) Surface smoothness of FCl alloy
by SEM image and (b) Structure by XRD pattern 3.2 Crystallization
Behavior of the Wires
The continuous scanning of the amorphous wires at the heating rate
of 20°C/min is shown by DSC thermograms (Figure 3). Two peaks in
FCl and single peak in FC4 signify the multi stage and single stage
crystallization of those alloys, respectively. As shown in Figure 3
and Table I, the onset and peak temperatures of FC 1 are lower than
that of FC4, representing higher thermal stability of FC4 alloy
compared to FCl alloy. The crystalline phases can be examined by
XRD patterns of the annealed wires (Figure 4). The first and second
crystallization peaks of FCl alloy correspond to α-FeCo phase and
FeB, CoB and CoSi phases, determined after annealing at 550 and
585°C, respectively. The Nb addition stabilizes ct-FeCo phases and,
therefore, Fe2Nb phases are predominant and no borides and/or
suicides phases are observed in FC4 alloy annealed at 630°C. The
crystalline size of the phases is measured by the broadening of the
X-ray diffraction patterns using Scherrer equation [11], D= 0.9λ/β
cos Θ, where D is crystallite size, λ is wavelength of incident
radiation (0.1540 nm), β is full width at half maximum (Table II).
The crystallite sizes of ct-FeCo phases decrease with the effect of
Nb addition. The element Nb acts as a growth inhibitor, resulting
in finer nanocrystallites in the amorphous matrix [12].
11
Figure 3 DSC thermograms of the amorphous wires at heating rate of
20°C/min
Table 1. Onset, Peak, End Temperatures and Activation Energy for
Crystallization of FC 1& FC4 Wires
Alloy Name
FC4
278 388 458
Table II. Crystallite Sizes (nm) of Various Phases Formed in FC1,
FC4 Alloys after Annealing Alloy name & Annealing
Temperature
FCl&ann. at 550°C FC1 &ann. at585°C FC4 & ann. at 575°C
FC4 & ann. at 630°C
Fe-Co Phase (nm)
12
The apparent activation energy of crystallization (Ea) for each
observed crystallization step can be determined by Kissinger's
relationship between the exothermic peak temperature (Tp) and the
heating rate (h) [13], described as the equation (1)
h E I n — - = — + constant (1)
T„2 RT„ where R is a gas constant.
According to equation (1), the plotting of ln(h/T p 2) as a
function of 1/TP yields a straight line
and activation energy (Ea) is determined from the slope-E JR) of
the lines (Table I). The activation energy of solid state reactions
is spent for overcoming and lowering of the activation barrier due
to rearrangements of atoms [14]. It results in the formation of
nuclei and their growth during crystallization. Therefore, the
energy calculated in these experiments, is determined for both the
lowering of the potential activation barrier and overcoming the
barrier. The activation energy of crystallization increases during
addition of Nb, indicating the role of a significant fraction of
the atoms in the structural reorganization. It causes the formation
of stable Fe2Nb phase in FC4 alloy, and improves the thermal
stability of the alloy.
3.3 Electron Transport Properties of Wires during Annealing
The order-disorder change and nanocrystallization behaviour in
amorphous alloys can be evaluated by electrical resistivity
measurement. Figure 5 represents the variation of normalized
resistivity of amorphous wires during isochronal annealing at the
heating rate of 10°C/min. In all the alloys, the resistivity
initially increases with temperature and then suddenly drops at the
Txi. It is attributed to the transformation of largely disordered
metastable amorphous state to the ordered crystalline state [15].
Therefore, Txi is the first crystallization temperature. The Txl
shifts to higher temperature with the addition of Nb in FC4 alloy.
The second crystallization (Τχ2) also occurs at higher temperatures
for FC4 alloy than FC 1 alloy. After the completion of second
crystallization, the growing nanophase particles lead to grain
boundary scattering and consequent increase in resistivity.
Fig. 5 Electrical resistivity measurements of amorphous wires at
the heating rate of 10°C/min
13
3.4 Variation of Magnetic Moments with Temperature
The thermal variation of magnetic moments of the as-cast wires was
measured to study the effect of crystallization process on Curie
temperature (Figure 6). The initial sharp drop is associated with
the curie temperature of the amorphous matrix. The curie
temperature is highest in FCl alloy and decreases with the addition
of Nb in other alloy. At Tc, the sudden drop of magnetic moments in
FCl and FC4 may be attributed to the rapid transition of
ferro-to-para magnetism, however, the continuous reduction of
magnetic moments in FC4 alloy is due to sluggish transition from
ferromagnetism to paramagnetism. The magnetic moment of FCl alloy
suddenly reverts back to initial stage, indicating ferromagnetic
coupling is largely present in that alloy and it is rapidly
changing its mode.
Figure 6. Change of magnetic moments as a function of
temperature
3.5 GM1 Properties of the Alloys
The composition dependence of GMI ratio is explained for the
as-received alloys in Figure 7. The curves show the single-peak GMI
characteristics behaviour and are sharpen with the addition of Nb
in the master alloy of Fe39Co39SisBi4 (FCl) for FC4 alloy. The
improvement of GMI properties is due to the change in skin depth of
the wire after addition of Nb [7, 16]. The skin effect of amorphous
wire is also dependent on ac frequency and amplitude of ac driving
field. Initially, the GMI ratio increases and then it decreases
above 400 kHz frequency (Figure 8a). The dependency of amplitude of
driving field on GMI effect also follows similar trend. The maximum
GMI ratio is observed at 10mA driving field (Figure 8b). The GMI
response is at highest when frequency and driving field are in the
range of 300 kHz < f < 500 kHz and 8 < I^ < 12,
respectively. It is noted that the addition of Nb is more effective
to enhance magnetic impedance in all ranges of frequency and ac
driving field. On the other hand, the skin depth is determined by
the circular permeability that is strongly frequency dependent
[16]. It causes the rise of GMI ratio to maximum range and follows
a decrease with an increasing frequency within the range from 100
kHz to 10MHz. Therefore, the enhanced GMI effect is a direct
consequence of the higher mobility of domain walls as correlated
with transverse permeability and transverse magnetic anisotropy
which are induced on the wire's surface during the rapid quenching
of wire in the water [17].
14
Figure 7. GMI ratio of amorphous wires at frequency of 400 kHz and
field amplitude of 10mA
Figure 8. GMI Ratio of different wires as a function of (a)
frequency at the driving field of amplitude 10mA and (b) amplitude
of driving field at frequency of 400 kHz.
Conclusions
It is concluded that crystallization kinetics and GMI properties of
amorphous wires are greatly changed with the addition of Nb in the
master alloy of Fe39Co39SisBi4 as follows:
1) Crystallization of FC1 alloy takes place in two stages, while
FC4 follows single stage crystallization.
2) Thermal stability and activation energy of crystallization
increases in FC4 alloy due to the formation of stable phases like
Fe2Nb.
3) The Nb addition does not affect much in curie temperature of FC4
alloy than that of FC1 alloy.
4) The GMI ratio is 18% for FC4 alloy and 10% for FC1 alloy. 5) The
GMI response of all alloys is at the highest when frequency and
driving field are in
the range of 300 kHz < f < 500 kHz and 8 < \x< 12,
respectively.
Acknowledgement
We express our thanks to the Director, National Metallurgical
Laboratory (CSIR), Jamshedpur, for giving us permission to publish
the work. The work is a part of the CSIR network project on
"Nanostructured Advanced Materials" (NWP-051).
15
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Intermetallics, 14 (2006), 1066-1068.
9. S. Li, S. Bai, H. Zhang, K. Chen, J. Xiao, "Effects of Nb and C
additions on the crystallization behavior, microstructure and
magnetic properties of B-rich nanocrystalline Nd-Fe-B ribbons", J.
Alloys Compd., 470 (2009), 141.
10. S.W. Du, R.V. Ramanujan, "Crystallization and magnetic
properties of Fe40Ni38B18Mo4 amorphous alloy", J. Non-Cryst.
Solids, 351 (2005), 3105.
U.M.P. Klug, L.F. Alexanader, X-ray Diffraction Procedures for Poly
crystalline and Amorphous Materials, (John Wiley & Sons, New
York, 1974) 634.
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M. E. McHenry, Vladimir Keylin, and Joe Huth," Increased induction
in FeCo-based nanocomposite materials with reduced early transition
metal growth inhibitors", J. of App. Phy, 107 (2010) 09A316.
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Analysis", Anal. Chem., 29 (1957), 1702.
14. D. .M. Minió, A. Gavrilovic, P. Angerer, D.G. Minie, A.
Marioic, " Thermal stability and crystallization of
FegçisNiisSisjBsCos amorphous alloy", J. of Alloys and Comp., 482
(2009), 502-507.
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R. Hasegawa (Eds.), Plenum Press, New York, 1977) 327.
16. M. Vazquez, J. of Mag. and Magn. Mater., "Giant
magneto-impedance in soft magnetic "Wires"", 226-230 (2001),
693-699.
17. N.D. Tho, N. Chau, SC. Yu, H.B. Lee, N.D. The, N.Q. Hoa, "A
systematic study of giant magnetoimpedance of Cr-substituted Fe(73
5-K)CrxSii35B9Nb3Aui (x=l, 2, 3, 4, 5) alloys", J. of Magn. and
Magn. Mater., 304 (2006), e871-e873
16
Fe-BASED AMORPHOUS-NANOCRYSTALLINE THERMAL SPRAY COATINGS
B. Movahedi'.M.H. Enavati2
'Faculty of Advanced Sciences and Technologies; University of
Isfahan; Isfahan, Iran department of Materials Engineering; Isfahan
University of Technology;
Isfahan, 84156-83111, Iran
Abstract
In this work, a new composition of Fe-15Cr-4Mo-5P-4B-lC-lSi (wt.%)
amorphous powder was produced by mechanical alloying of elemental
powder mixture. Thermal spraying of amorphous powder was done by
high velocity oxy fuel spraying technique at various spraying
conditions to obtain the desirable amorphous and nanocrystalline
coatings. It was found that a-Fe based supersaturated solid
solution is first formed during mechanical alloying which
transforms to amorphous structure at longer milling times. The
crystallization kinetic parameters suggest that the crystallization
mechanism is dominantly governed by a three-dimensional diffusion-
controlled growth. The crystallization of amorphous structure
occurs in one single stage. By carefully controlling the spraying
parameters and proper selection of powder composition, the desired
microstructure with different fraction of amorphous and
nanocrystalline phases and therefore with different properties
could be obtained.
Introduction
Amorphous metallic alloys have been of interest not only for
fundamental studies, but also for potential applications for over
40 years. Fe-based amorphous alloys are perhaps the most important
system for possible applications because of the low cost of iron,
and the relatively high strength and hardness of Fe-based amorphous
alloys [1]. The formation of amorphous phase by mechanical alloying
(MA) process depends on the energy provided by the milling machine
and thermodynamic properties of the alloy system. There are two
rules for the formation of amorphous alloy by MA in an A-B binary
system: (1) a large negative heat of mixing, AHm¡x, between the
elemental constituents and (2) a large asymmetry in the diffusion
coefficients of the constituents. An amorphous phase is kinetically
obtained only if the amorphization reaction is much faster than
that for the crystalline phases [2]. Synthesizing amorphous and/or
nanocrystalline coatings on metal substrates can be utilized to
improve surface performance such as wear and corrosion resistance.
Thermal spraying process is one of the techniques to deposit
amorphous coatings on surfaces, where the amorphous structure is
retained due to the sufficiently rapid cooling that inhibits
long-range diffusion and crystallization [3]. A number of
researchers have investigated the use of air plasma spraying (APS)
and high velocity oxy fuel (HVOF) to deposit alloys, which are
capable of solidifying as metallic glasses [4]. In this work a new
composition of Fe-Cr-Mo-B-P-Si-C amorphous powder was first
prepared by mechanical alloying of elemental powder mixtures. In
next step this amorphous powder was sprayed by high velocity oxy
fuel (HVOF) spraying techniques to obtain amorphous and
17
nanocrystalline coatings. The microstructure and tribological
behavior of coatings were investigated in details by X-ray
diffractometry (XRD), scanning electron microscopy (SEM),
transmission electron microscopy (TEM), differential scanning
calorimetry (DSC) and wear tests.
Experimental The elemental powders were blended to give a nominal
composition of 70Fe-15Cr-4Mo-5P-lC- 1SÍ-4B (wt.%). The purity and
mean particle size of as-received powders are given in Table 1. Red
phosphorus had an amorphous structure while the rest of
constituents were crystalline.
Table 1. Purity and mean particle size of as-received powders.
Element
Iron Chromium Molybdenum Borne Graphite Red Phosphorous
Silicon
Mean particle size (μπι)
99.00% 99.90% 99.00% 98.00% 99.99% 99.00% 99.90%
Mechanical alloying was performed in a high-energy planetary ball
mill (Retch PM100) in argon atmosphere using hardened chromium
steel vial and balls (Φ=20 mm). The ball-to-powder weight ratio was
10:1 and the rotation speed of the main disc was 280 rpm. The MA
was done nominally at room temperature although the temperature of
the vial increased to around 50°C during MA. The milling was
interrupted at different selected times and a small amount of
powder was taken out of the vial for further analysis. The MA
powder was sprayed on a carbon steel substrate (50 by 50 by 5mm)
using HVOF (Metallisation Met JET II) system with different
parameters as shown in Table 2.
Table 2. HVOF spraying parameters
Parameters
Oxygen gas flow rate (SLPM) Fuel (Kerosene) flow rate (SLPM)
Fuel/Oxygen (Vol%) Powder feed rate (g min-1) Spray distance (mm)
Scanning velocity (mm s-1) Deposit thickness (μπι) Nozzle length
(mm) Compress air cooling
Microstructure
18
X-ray diffraction (XRD) was performed to study the structural
evolution of powders during the ball milling process. Differential
scanning calorimetry (DSC) with a constant heating rate of 20 K/min
under flowing argon gas (99.999%) was used to study the
crystallization behavior of amorphous powder. The morphology and
cross-sectional microstructure of powder particles after different
milling times were investigated by scanning electron microscopy
(SEM). High resolution transmission electron microscopy (HRTEM) of
powder particles was carried out using a Jeol-JEM-2010 TEM at an
accelerating voltage of 200 kV and resolution of 0.19nm.
Results and discussion Development of amorphous structure
Figure 1 shows the XRD patterns of powder mixture as a function of
milling time. As-received powder mixture shows sharp crystalline
peaks of elemental Fe, Cr, Mo, B, C and Si. Red Phosphorus is
absent on XRD pattern because it's amorphous nature. As milling
progresses, the XRD peaks of the elemental constituents are
broadened with a corresponding decrease in their intensities. These
effects are caused by a continuous decrease in effective
crystalline size and an increase of the atomic level strain, as a
result of the induced-plastic deformation during MA [5]. On
continued milling a broad peak was developed on the XRD pattern,
owing to the formation of an amorphous phase. A fully amorphous
structure was obtained after 80 h of milling time.
Figure 1. XRD patterns of Fe-Cr-Mo-B-P-Si-C powder mixture as
received and after different milling times.
Microstructural observations of powder HRTEM images, selected-area
diffraction patterns (SADP) and fast Fourier transform (FFT) images
of powders milled for 15 h (Figure 2a) confirmed the formation of a
nanocrystalline structures. After 40 h of milling time amorphous
and nanocrystalline phases co-existed in the milled powders. Figure
2b shows that most amorphous phase are developed at the edge of
powder particles indicating that the amorphization reaction starts
at edge of particles and progress into the internal regions as MA
proceeds [6]. Figure 2c is the HRTEM image and SADP of powder after
80 h of milling time, showing a fully amorphous
microstructure.
19
Figure 2. HRTEM micrographs, SADP and FFT patterns of
Fe-Cr-Mo-B-P-Si-C amorphous powder after different milling
times.
The structure of coatings Figure 3 illustrates the XRD patterns of
mechanically alloyed Fe-Cr-Mo-P-B-C-Si feedstock powder and the
as-sprayed HVOF coatings. The XRD pattern of HVOF-G1 coating in
Figure 3 has a halo characteristic indicating that this coating has
an amorphous structure similar to feedstock MA powder. However in
HVOF-G2 there is an emergent crystalline peak on the top of the
amorphous hub suggesting that this coating is a mixture of
amorphous and crystalline phases. Structure of HVOF-G3 coating
mainly consists of crystalline phases such as a-Fe, Fe23(C, B)6 and
Fe5C2.
20
Figure 3. XRD patterns of mechanically alloyed feedstock powder and
HVOF coatings.
It is inferred from the XRD results that a range of microstructures
from fully amorphous to fully crystalline can be obtained by
adjusting of HVOF parameters (see Table 2). The difference in the
fraction of amorphous phase is related to the amount of cooling
rate and remelting of individual particles in HVOF flame at various
fuel/oxygen ratios. By increasing the flame temperature the powder
particles are completely remelted in flame and then rapidly
solidified and quenched on the cold substrate forming an amorphous
structure.
Figure 4. HRTEM micrograph and SADP of fully amorphous HVOF-G 1
coating.
HRTEM image (Figure 4) confirms that HVOF-G 1 coating is completely
amorphous. As shown in Figure 5 the HVOF-G2 coating consists of
amorphous phase and nanocrystalline grains of 5- 30 nm. In this
case the fuel/oxygen ratio is moderate (HVOF-G2) therefore, this
duplex
21
microstructure can be explained by quenching of semi-molten
particles when impinged to the cold substrate. A nanocrystalline
structure with equiaxed grains was obtained in case of HVOF-G3
coating (Figure 6). In this condition the fuel/oxygen ratio has a
minimum value and the HVOF flame temperature is the lowest
therefore, the most of the individual powder particles were
unmelted and crystallized inside the HVOF flame. Moreover, the
cooling rate was sufficiently high to avoid grain coarsening
yielding a nanocrystalline structure.
Figure 5. a) TEM and b) HRTEM micrographs, SADP and FFT of
amorphous-nanocrystalline HVOF-G2 coatings.
Figure 6. TEM and HRTEM micrographs and SADP of fully
nanocrystalline HVOF-G3 coating.
Thermal behavior Figure 7 shows DSC traces of as-milled powder and
HVOF coatings. As seen the crystallization of MA powder and
coatings occurs in a single stage around 560-580°C. The supercooled
liquid region, ΔΤΧ, defined by the difference between the glass
transition temperature (Tg) and the onset
22
temperature of crystallization (Tx), is as large as 69 °C (Table
3). It is suggested that a large ΔΤΧ generally represents a high
glass forming ability (GFA) in the amorphous alloys [7].
Figure 7. DSC traces of mechanically alloyed feedstock powder and
HVOF coatings.
Table 3. Crystallization characteristics of Fe-Cr-Mo-P-B-C-Si
powder and coatings. Microstructure
Mechanical alloying powder
Amorphous (HVOF-G1) Amorphous-
100
100
44
2.1
The Avrami exponents for different temperatures range from 2.34 to
3.32, which imply that the crystallization mechanism depends on
temperature during non-isothermal annealing. At 570 and 572CC, the
values of (n) are 3.32 and 3, respectively that is typical for
interface controlled two dimensional growth of nuclei with
decreasing nucleation rate. The Avrami exponent values decrease to
2.54 when the temperature increases to 576°C, suggesting that the
growth mechanism changes to the volume diffusion controlled three
dimensional growth of nuclei with constant nucleation rate. The
high value of activation energies of crystallization Ea
(386.04kJ/mol) indicates that a lot of atoms participate in an
elementary act of structural reorganization so that the atomic
diffusion in Fe-Cr-Mc—B-P-Si-C system is difficult, especially at
low temperature demonstrating the MA amorphous Fe-Cr-Mo-B-P-Si-C
powder exhibits high glass forming ability and thermal stability
[8, 9].
23
Conclusions A fully amorphous structure was obtained by mechanical
alloying of 70Fe-15Cr-4Mo-5P-lC- 1SÍ-4B (wt.%) powder mixture.
Amorphization reaction appeared to start at edge of powder
particles and progresses into the internal regions as mechanical
alloying proceeds. The results also indicated that this alloy
system has a high tendency to form amorphous structure by
mechanical alloying with high GFA and thermal stability. The
significant variation of the local Avrami exponent and local
activation energy for crystallization demonstrated that the
crystallization kinetics varies at different stages. The
crystallization process is mainly governed by three-dimensional
diffusion-controlled growth of nuclei. The large ΔΤΧ and activation
energy of crystallization indicate the high thermal stability of
this amorphous alloy produced by high energy mechanical alloying.
HVOF spraying of mechanically alloyed amorphous Fe-Cr-Mo-P-B-C-Si
powder were employed to obtain amorphous and nanocrystalline
coatings. It was showed that thermal spraying techniques are able
to prepare a wide range of microstructure from amorphous to
nanocrystalline in Fe-Cr-Mo-P-B-C-Si alloy system. At low flame
temperature a partial or full crystallized coating was obtained
while spraying at higher flame temperatures led to a fully
amorphous structure.
References 1. A.L. Greer, K.L. Rutherford, and I.M. Hutchings,
"Wear Resistance of Amorphous Alloys
and Related materials," International Materials Review, 47 (2002),
87-112.
2. C. Suryanarayana, "Mechanical Alloying and Milling," Progress in
Materials Science, 46 (2001), 1-184
3. Y. Wu, P. Lin, G. Xie, J. Hu, and M. Cao, "Formation of
Amorphous and Nanocrystalline Phases in High Velocity Oxy-Fuel
Thermally Sprayed Fe-Cr-Si-B-Mn Alloy," Materials Science and
Engineering A, 430 (2006), 34-39
4. K. Kishitake, H. Era, and F. Otsubo, "Characterization of Plasma
Sprayed Fe-Cr-Mo-(C, B) Amorphous Coatings," Journal of Thermal
Spray Technology, 5 (1996), 145-153
5. M.S. El-Eskandarany, W. Zhang, and A. Inoue, "Mechanically
Induced Crystalline-Glassy Phase Transformations of Mechanically
Alloyed TaZrAlNiCu Multicomponent Alloy Powders," Journal of Alloys
and Compounds, 350 (2003), 222-231
6. P. Schumacher, M.H. Enayati, and B. Cantor, "Amorphization
Kinetics of Ni60Nb40 During Mechanical Alloying," Journal of
Metastable and Nanocrystalline Materials, 2-6 (1999), 351-356
7. X. Wu, and Y. Hong, "Fe-based Thick Amorphous-Alloy Coating by
Laser Cladding," Surface and Coating Technology, 141 (2001),
141-144.
8. S.J. Pang, T. Zhang, K. Asami, and A. Inoue, "Synthesis of
Fe-Cr-Mo-C-B-P Bulk Metal Glasses with High Corrosion Resistance,"
Acta Materialia, 59 (2002), 489-497
9. B. Movahedi, "Microstructural and Tribological Evaluation of
Novel Fe-based Amorphous- Nanocrystalline Thermal Spray Coatings"
(Ph.D. thesis, Isfahan University of Technology, 2010),
85-162.
24
Enhanced photocatalytic activity of modified Ti02 for
degradation of CH20 in aqueous suspension Haixia Tonga'b*, Li
Zhaoc, Dan Li"·b and Xiongfei Zhang'
* Chemical and Biologic Engineering Institute, Changsha University
of Science and Technology, Hunan
Province Key Laborator of Materials Protection for Electric Power
and Transportation,
Changsha 410076, Hunan, China
China c College of Chemistry and Chemical Engineering, Nanjing
University, Nanjing 210093, Jiangsu, China
ABSTRACT
Butyltitanate, ethanol and glacial acetic acid were chosen as
titanium source, solvent and
chelating agent respectively via a sol - gel method combined
impregnation method to prepare
N, Fe co-doped and W 0 3 compounded photocatalyst TiC>2 powder.
The synthesized products
were characterized by X-ray diffraction (XRD), Diffuse reflectance
UV-Vis spectra
( UV-DRS ) , and scanning electron microscopy ( SEM ) . The
catalytic activity was
investigated employing photocatalytic degradation of formaldehyde.
The results show that the
degradation rate is 77.61% in 180 min under UV light irradiation
when the concentration of N
is fixed on, and the optimum proportioning ratio of n (Fe): n (W):
n (Ti) is 0.5:2:100.
KEY WORDS : N-Fe Co-doping; W 0 3 compounded; photocatalysis;
formaldehyde
»Corresponding author: Ph. D; Tel: +86- 731-85258733, Fax:
+86-731-85258733 E-mail:
tonghaixia@,l 26.com
1. Introduction
Nowadays, because of the extensive using of interior decorative
materials and household
chemicals, more and more attention is paid to the research of
indoor air pollution. The Volatile
Organic Compounds (VOCs) are a class of major indoor air
pollutionsfl]. Formaldehyde is
one of the typical VOCs aldehydes among the indoor air pollutions
at present, and it has been
recommended as one of the model compounds to test the performance
of air filtration
equipments by the U.S. ASHRAE (American Society of Heating,
Refrigerating, and
Air-Conditioning Engineers )[2]. Therefore, how to eliminate indoor
formaldehyde effectively
has become an increasing hot research topic.
Recently, the photocatalytic oxidation technology has been applied
in air purification,
and the degradation of VOCs in the air also has been caused an
extensive research[3-8].
Among of various photocatalysts including T1O2, ZnÛ2, W0 3 , CdS,
ZnS, SrTi03, and Fe203,
T1O2 has received much research interest due to its chemical
stability, nontoxicity and high
photocatalytic activity.
However, as a wide bandgap semiconductor (3.2 eV anatase), Ti02can
absorb only the
UV light of solar energy, which limits their practical application.
At the present time, many
researches are carried out in order to substitute the UV light by
the sunlight or visible light. This will reduce
the cost of the photodegradation process especially for industrial
scale applications. Using dopants that can
be incorporated in T1O2 lattice is one of the methods used to reach
this goal[9-17]. Results show that doping
T1O2 with transition metals and (or) non-metal elements (S, N, C
and F) increases its photocatalytic
activity [13-17]. It is reported that doped ions can enhance the
intensity of absorption in the UV-vis light
region and make a red shift in the band gap transition of the doped
Ti02 samples. For example in the case
of Fe-Ti02 sample, Fe ions can have two roles: they can act as a
photo-generated hole and a
photo-generated electron trap and reduce the hole-electron
recombination and also they can serve as a
mediator of the transfer of interfacial charge [13,17]. However,
there is a controversy on the effect of
metals ions on the photocatalytic activity of Ti02. Other authors
show a decrease of photocatalytic activity
of the doped catalysts [18-21]. The amount of the metalions that
can be incorporated inTiC>2 lattice is also a
controversial matter [13,15,16].
26
Some studies have also shown that semiconductor compounding is
beneficial to the
separation of the photo-electrons and holes, which can enhance the
photocatalytic efficiency.
For example, Zhang Qi et. al. [22] reported that W03 thin films
sputtered on Ti02 can improve
the speed of the photocatalytic degradation of méthylène blue. In
our previous works[23] a
suitable amount WO3 compounding can improve the photocatalytic
activity of Ti02 in
splitting water and the optimum concentration of compounded WO3 is
2 %. However there is
no report on T1O2 modified by Fe, N and WO3 at the same time.
In this paper, Ti02 photocatalysts doped with N, Fe and compounded
by W03 are
composited, and used for photo degradation of formaldehyde solution
under the UV-light
irradiation. The Fe element is providied by Fe (NO3) 3 · 9H20,
which is cheap, non-toxic and
simple doping process, and N element is providied by ammonia, which
is also cheap, simple
doping process and easy to be practically used. Researches of the
catalysts used for photo
degradation of formaldehyde under the Vis-light irradiation are in
progress.
2. Experimental section
2.1 Catalyst preparation
(1) 17mL butyl titanate was added to 40mL anhydrous ethyl alcohol
drop by drop and
then stirred continually for 30 min with magnetic stirrer. A yellow
transparent solution was
obtained at room temperature, and was titled A;
(2) lOmL glacial acetic acid was added to 5mL distilled water, and
then shaked the
mixture up, and added to 40mL ethanol. The solution B was
obtained;
(3) Solution A was dropwise added to solution B under vigorous
stirring at room
temperature, and then adjusted pH value to 1 ~ 2 with concentrated
hydrochloric acid. When
the color of the solution turned to light yellow and continued to
stir for half an hour.
(4) After aging at room temperature for 24 h, the obtained sol
solution was dried at
40°Cuntil a dry gel was gotten. This gel was grinded and calcined
in air for 4 h at 500 °C,
with a constant heating rate of 1 "Cmkf'. After grinded for 1 h,
crystalline Ti02 particles was
obtained, and denoted as: T (0).
27
In step (3), before the pH adjustment, 10 drops of stronger ammonia
water were added
and other steps were unchanged. At last the N doped T1O2 was
obtained, and denoted as: T
(N);different amounts of Fe (NO3) 3 · 9H2O solid was dissolved in
the appropriate anhydrous
ethanol, and added into the mixture before the ammonia water, and
other steps were
unchanged. At last a series of N, Fe co-doped Ti02 catalysts:
N-0.1% Fe-Ti02, N-O.5% Fe-
T1O2, N-0.7% Fe- Ti02, N-1.0% Fe- Ti02 were obtained, and denoted
as: T (NF1), T (NF2 ),
T (NF3), T (NF4), respectively.
The prepared N-0.5%Fe-TiO2 (T (NF2 )) powder was immersed in
different
concentrations of APT (Ammonium paratungstate) solution, and
grinded for half an hour, then
dried in an infrared oven for 2 h at 100°C. After grinded for half
an hour, the mixture was
calcined at 500 °C for 4 h, and then cooled to room temperature;
grinded and the powders of
N-0.5% Fe-xW(V Ti02 catalyst were obtained, and the value of x was
0.5%, 1%, 2%, 4%,
6% respectively, the catalyst powders were denoted as: T(NFW1),
T(NFW2), T (NFW3),
T(NFW4), T (NFW5) respectively.
2.2. Characterization of photocatalysts
X-ray diffraction analysis (XRD) was used to check the coexistence
of different crystal
phases of the catalyst by a HATCHI D/max2250 powder X-ray
diffractometer. The diffraction
profiles were recorded with Cu Και radiation (0.154056 nm) over a
2Θ range of 10 to 90 . A
plumbaginous counter with monochromator was used. The X -ray tube
was operated at 40 kV
and 300 mA.
Diffuse reflectance UV-Vis spectra ( UV-DRS ) measurements were
carried out on a
Beijing Purkinje TU-1901 UV/Vis spectrophotometer equipped by a
diffuse reflectance
accessory with an IS 19-1 integrating sphere, and BaS04 powder was
used as reference.
Surface morphology of the T1O2 catalysts were examined by scanning
electron
microscope (Japan JSM-5600LV). Small pieces of the prepared
photocatalysts were stuck on
stubs using double-sided tape. Before the sam-ples were analyzed,
they were sputtered with
a layer of gold film to prevent the occurrence of charging
effect.
28
2.3. Photocatalytic degradation of formaldehyde
The photocatalytic activity of the T1O2 catalysts were studied by
degradation of formaldehyde as a
target pollutant.The experiments were carried out in a 250 ml
cylindrical glass reactor inside equipped with
an ultraviolet (UV) lamp (365 nm, 250 W) using 250 ml formaldehyde
solution with an initial
concentration of 30 μg / mL and 0.5 g catalyst. Before the
photocatalytic degradation, the
suspension was magnetically stirred in the dark for 30min to
establish a formaldehyde
adsorption /desorption equilibrium. 5 ml solution was collected
from the suspension and was
immediately centriftiged at 4000 r/ min for 10 min. The
concentration of formaldehyde was
determined by spectroscopic analysis at 270 nm using a TU-1900 UV
spectrometer (Beijing
purkinje general instrument Co.Ltd., China). The corresponding
formaldehyde degradation rate was
calculated according to the following equation:
7 7 = ^ ^ x 1 0 0 % (1)
where A0 = the initial concentration pollutant, A = the
concentration of model pollutant at
experimental time t.
3.1. Crystal structure and morphology of Ti0 2 catalysts
XRD analysis was carried out to coifirm the polymorphs and
crystalline phases of T1O2
catalysts. The XRD patterns for the modified T1O2 catalysts are
shown in Fig.l. From the
patterns some characteristic peaks for T1O2 can be observed, and
the diffraction peaks locating
at 2e(deg.)=25.26 , 37.8 and 48.0° could assign to the planes of
(101), (111) and (200),
respectively, which all match well with the anatase Ti02. After N
doping, compared with pure
T1O2, the lattice parameter in c-axils has some extent increase
(c=9.4553 and 9.4994 nm for
T(0) and T(N), respectively), suggesting the N had inserted into
the crystal lattice of Ti02.
The patterns of T(NF2) before calcined show that the diffraction
peaks locating at 26(deg.)=
23.24 and 33.8°, which all match well with the NH4CI. However,
after calcined the peaks of
NH4CI phase disappeared, which also indicated that after calcined
at 500 °C for 4 h, N had
29
entered into T1O2 lattice.
However, before and after of calcined, T (NF2) has no peaks of
Fe203 because of lower
iron ions added .
XRD analysis does not present any characteristic diffraction peak
of W03 phase or other
tungsten oxide phase in the pattern of T (NFW3), and the pattern of
T (NFW6) shows the
characteristic diffraction peaks of W03. W03 can be well dispersed
on the Ti02 phase when
its concentration is less than 2%. WO3 is gathered and crystallized
when the concentration of
compounding WO3 is more than 2%[22], and the characteristic
diffraction peaks can be
presented on the XRD pattern.
SEM images of the catalysts at magnification of 30,000 and 5,000
times are shown in
Fig.2. and Fig.3. respectively. From the SEM images the catalysts
powder agglomerates
significantly. Fig.2 shows that T (NF1) is sculptured "pattern"
compared with other powders,
while the T (NF4) in the "pattern" is less clear than that of T
(NF1). Viewed the SEM of 5,000
times, the spherical morphology of T (NF1) powder is obvious;
viewed from Fig.3 T (0) and T
(N) agglomerate significantly more serious than that of Fe-doping
and WO3 compounding
T1O2 powders, which indicates that Fe3+ doping can reduce the
particle size. How about the
role of WO3 compounding on the particle size needs further
verification.
3.2 UV-DRS analysis
The DRS results of the catalysts with different Fe3+ contents are
shown in Fig.4. The
experimental results indicated that the undoped T1O2 powder (T(0))
shows strong
photoabsorption only at wavelengths shorter than 400 nm, which is
the characteristic
absorption of the charge transfer of O 2pTi 3d resulted from the
charge transfer from
oxygen atom coordinated with titanium to the empty orbit of the
center titanium atom [24].
While Fe3+ and N-doped T1O2 nanoparticles show photoabsorption in
visible region and the
absorption edge shifts to a longer wavelength. With the same
N-doped contents, the shift of
the reflectance spectrum is due to increasing Fe3+contents. This
indicates a decrease in the
band gap of Ti02.
30
Fig.5 shows the rfiectance spectra of T(0), T(NF2) and T(NF2) with
different WO3
compounding . Obviously, the absorption of T(NF2) with different
W03 compounding larger
than that of T(0), T(NF2) nanoparticles in the visible rigion. With
the increase of
compounding W03 concentration, T(NF2) catalysts exhibit wider
optical absorption, because
coupled WO3 can introduce an impurity level between the valence and
conduction band of
T1O2 and decrease its band gap[25].
3.3 Photocatalytic activity investigation
To evaluate the photocatalytic activity of the N and Fe-doped Ti02
photocatalysts,
degradation of CH20 solution were run under UV irradiation which
are graphically illustrated
in Fig.6. From thefigure, N and Fe-doped Ti02 exhibits much greater
activity than that of
pure Ti02. Fig.7 reflects the photocatalytic activity of Ti02
photocatalysts with different Fe
contents and same N content. The experiment results indicate that
the photocatalytic activity
of T1O2 photocatalysts with same N content can be improved by
doping an appropriate
content of Fe. When the Fe content is 0.5 mol %, the photocatalyst
exhibits higher
photocatalytic activity. If the Fe content continuously increases,
the photocatalytic activity
begins to fall down inversely. The degradation rates of T(N), T
(NF1), T (NF2), T (NF3) and
T (NF4) are 30.99%, 37.635%, 60.57% 54.54% and 51.885%
respectively.
The photocatalytic activities of T (NF2) photocatalysts with
different WO3 compounding
are shown in Fig.6, and the relationship between W03 concentration
and the CH20 degradation rate
is shown in Fig.8. During the degradation process for 3 h, the
degradation rate for T (NFW1), T
(NFW2), T ( NFW3), T (NFW4) and T (NFW5) are 62.055%, 70.47% 77.61%
58.785% and
40.38% respectively. It is easy to conclude that the optimize
concentration of W03
compounding is 2 mol %, which is consistent with the reported
result [22, 23].
During the calcination process N and Ti02 can form Ti02.xNx, and a
new energy band
which narrows the band gap of Ti02 has been introduced in. The
doping band introduced from
the mixture of substitutional N 2p and O 2p orbits is responsible
for the gap narrowing[26].
When it comes to the impact of doping Fe3+, it generally considers
that Fe3+ can capture
31
photoproduced electrons, and reduces the probability of electrons
and holes to recombination
and extends the average life expectancy of the holes, which is
beneficial to improve the
photocatalytic activity[27]. In addition, the study also shows that
Fe3+ can be e-captured, due
to the facts that the energy levels for Fe2+/Fe3+ and Ti3+/Ti4+ are
close, and thus the capturing
electro