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
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2nd International Symposium on High-Temperature Metallurgical Processing
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Magnesium Technology 2011
Recycling of Electronic Waste II, Proceedings of the Second Symposium
Sensors, Sampling and Simulation for Process Control
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Supplemental Proceedings: Volume 1: Materials Processing and Energy Materials
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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
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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.
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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.
References
1. E. Reddington, A. Sapienza, B. Gurau, R. Viswanathan, S. Sarangapani, E. S. Smotkin, and T. E. Mallouk, "Combinatorial Electrochemistry: A High Parallel, Optical Screening Method for Discovery of Better Electrocatalysis," Science, 280 (1998), 1735-1737.
2. R. X. Liu, and E. S. Smotkin, "Array membrane electrode assemblies for high throughput screening of direct methanol fuel cell anode catalysts," J. Electroanal. Chem., 535 (2002), 49-55.
3. Y. Lou, M. M. Maye, L. Han, J. Luo, and C.-J. Zhong, "Gold-platinum alloy nanoparticle assembly as catalyst for methanol electrooxidation," Chem. Commun., 5 (2001), 473-474.
4. V. R. Stamenkovic, B. S. Mun, M. Arenz, K. J. J. Mayrhofer, C. A. Lucas, G. Wang, P. N. Ross, N. M. Markovic, "Trends in electrocatalysis on extended and nanoscale Pt-bimetallic alloy surfaces," Nat. Mater., 6 (2007), 241-247.
5. M. Monta, Y. Iwanaga, and Y. Matsuda, "Anodic oxidation of methanol at a gold-modified platinum electrocatalyst prepared by RF sputtering on a glassy carbon support," Electrichim. Acta, 36 (1991), 947-951.
7
6. L. D. Burke, J. A. Collins, M. A. Horgan, L. M. Hurley, and A. P. O'Mullane, "The importance of the active states of surface atoms with regard to the electrocatalytic behaviour of metal electrodes in aqueous media," Electrochim. Acta, 45 (2000), 4127-4134.
7. H. M. Chen, H-. C. Peng, R-. S. Liu, S-. F. Hu, and H-. S. Sheu, "Morphology and Surface Plasma Changes of Au-Pt Bimetallic Nanoparticles," Nanosci. Nanotechnoi, 6 (2006), 1411-1415.
8. D. Mott, J. Luo, P. N. Njoki, Y. Lin, L. Wang, and C-. J. Zhong, "Synergistic activity of gold-platinum alloy nanoparticle catalysts," Catalysis Today, 122 (2007), 378-285.
9. P. H-. Fernández, S. Rojas, P. Ocón, J. L. Gómez de la Fuente, J. S. Fabián, J. Sanza,M. A. Pefla, F. J. G-. Garcia, P. Terreros, and J. L. G. Fierro, "Influence of the Preparation Route of Bimetallic Pt-Au Nanoparticle Electrocatalysts for the Oxygen Reduction Reaction," J. Phys. Chem. C, 111 (2007), 2913-2923.
10. J. K. Lee, J. Lee, J. Hanc, T-. H. Lime, Y-. E Sungd, and Y. Tak, "Influence of Au contents of AuPt anode catalyst on the performance of direct formic acid fuel cell," Electrochimica Ada, 53 (2008), 3474-3478.
11. M. Schrinner, S. Proch, Y. Mei, R. Kempe, N. Miyajima, and M. Ballauff, "Stable Bimetallic Gold-Platinum Nanoparticles Immobilized on Spherical Polyelectrolyte Brushes: Synthesis, Characterization, and Application for the Oxidation of Alcohols," Adv. Mater., 20 (2008), 1928-1933.
12. J. Luo, M. M. Maye, V. Petkov, N. N. Kariuki, L. Wang, P. Njoki, D. Mott, Y. Lin, and C-. J. Zhong, "Phase Properties of Carbon-Supported Gold-Platinum Nanoparticles with Different Bimetallic Compositions," Chem. Mater., 17 (2005), 3086-3091.
13. J. Luo, P. N. Njoki, Y. Lin, D. Mott, L. Wang, and C-. J. Zhong, "Characterization of Carbon-Supported AuPt Nanoparticles for Electrocatalytic Methanol Oxidation Reaction," Langmuir, 22 (2006), 2892-2898.
14. T. Nakamura, Y. Mochidzuki, and S. Sato, "Fabrication of gold nanoparticles in intense optical field by femtosecond laser irradiation of aqueous solution,"/ Mater. Res., 23 (2008), 968-974.
15. T. Nakamura, K. Takasaki, A. Ito, and S. Sato, "Fabrication of platinum particles by intense, femtosecond laser pulse irradiation of aqueous solution," Appl. Surf. Sei., 255 (2009), 9630- 9633.
16. S. L. Chin, and S. Lagacé, "Generation of H2, 02 , and H2O2 from water by the use of intense femtosecond laser pulses and the possibility of laser sterilization," Appl. Opt., 35 (1996), 907-911.
17. H, H, Huang, X. P. Ni, G. L. Loy, C. H. Chew, K. L. Tan, F. C. Loh, J. F. Deng, and G. Q. Xu, "Photochemical Formation of Silver Nanoparticles in Poly(N-vinylpyrrolidone)," Langmuir, 12 (1996), 909-912.
18. A. Henglein, "Colloidal Silver Nanoparticles: Photochemical Preparation and Interaction with O2, CCU, and Some Metal Ions," Chem. Mater, 10 (1998), 444-450.
19. T. Kempa, R. A. Farrer, M. Giersig, and J. T. Fourkas, "Photochemical Synthesis and Multiphoton Luminescence of Monodisperse Silver Nanocrystals," Plasmonics, 1 (2006), 45-51.
20. G. C. Bond, "The Electronic Structure of Platinum-Gold Alloy Particles," Platinum Metals Rev., 51 (2007), 63-68.
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
1. L.V. Panina, K. Mohri, "Giant magnetic field dependent impedance of amorphous FeCoSiB wire", Appl Phys Lett, 65 (1994), 1189-91.
2. T. Meydan, J Magn Magn Mater, "Application of amorphous materials to sensors", 133 (1995), 525-32.
3. J.E. Lenz, "A Review of Magnetic Sensors", Proc. IEEE, 78 (1990), 973-89. 4. M. H. Phan, H. X. Peng, "Giant magnetoimpedance materials-Fundamentals and
applications'" Prog, in Mat. Sc, 53 (2008), 323-420. 5. A.O. Olofinjana, J.H. Kern, H.A. Davies, "Effects of process variables on the multi-
strand casting of high strength sub-millimetre metallic glass wire", J. Mater. Proc. Tech., 155-156(2004), 1344-1349.
6. Z. Stoklosa, J. Rasek, P. Kwapulinski, G. Haneczok, G.Bzdura, J. Lelajtko, "Nanocrystallisation of amorphous alloys based on iron", Mat. Sc. andEngg C, 23 (2003), 4 9 - 5 3 .
7. D. Y. Liu, W.S. Sun, H.F. Zhang, Z.Q. Hu," Preparation, thermal stability and magnetic properties of Fe-Co-Ni-Zr-Mo-B bulk metallic glass" Intermetallics, 12 (2004), 1149— 1152.
8. W.S. Sun, T. Kulik, X.B. Liang, J. Ferenc, "Thermal stability and magnetic properties of Co-Fe-Hf-Ti-Mo-B bulk metallic glass", 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.
12. K. J. Miller, A. Leary, S. J. Kernion, A. Wise, D. E. Laughlin, 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.
13. H. E. Kissinger, "Reaction Kinetics in Differential Thermal 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.
15. W. Teoh, N. Teoh, S. Arajs, Amorphous Magnetism II, (R. Levy, 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.
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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.
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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
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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].
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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.
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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].
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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.
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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.
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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