TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings [1st ed.] 978-3-030-05860-9, 978-3-030-05861-6

Table of contents :
Front Matter ....Pages i-xxii
Front Matter ....Pages 1-1
Ab Initio Molecular Dynamics Study on the Dissolution of Interfacial Iron Oxides in Hot Compressive Bonding Combined with Experiments (Honglin Zhang, Mingyue Sun, Bin Xu, Dianzhong Li)....Pages 3-15
Effect of MgO Content on the Properties of Magnesia Fluxed Pellets (Yuzhu Zhang, Weixing Liu, Aimin Yang, Jie Li)....Pages 17-28
Numerical Simulation of Three-Phase Flow of Gas-Stirring Micro-phenomenon During Ladle Furnace Process (Libin Zhu, Wei Liu, Shfueng Yang, Jingshe Li, Feng Wang, Xueliang Zhang)....Pages 29-38
The Effect of pH and Temperature During Carbonation Process on Spent Die Cleaning Solution from Aluminium Extrusion Industry (Ahmed S. Aadli)....Pages 39-49
Improvement of Center Segregation in Continuously Cast Blooms by Convex Roll Soft Reduction (Liang Li, Xiao Zhao, Peng Lan, Zhanpeng Tie, Haiyan Tang, Jiaquan Zhang)....Pages 51-61
Effects of a Top-Down Flow on Gas–Solid Fluidization State in a Bubble Fluidized Bed (Xu Han, Liangying Wen, Shengyun Shi, Wenhuan Jiang, Meihuan Liu, Feng Lu)....Pages 63-76
Development of Bio-treated Oil Palm Fiber Reinforced Kaolin Matrix Composites for Building Bricks Application (Muideen Adebayo Bodude, Olasunkanmi B. Adegbuyi, Ruth Nkiruka Nnaji)....Pages 77-91
Effect of Roll Surface Profile on Thermal-Mechanical Behavior of Continuously Cast Bloom in Soft Reduction Process (Liang Li, Xiao Zhao, Peng Lan, Zhanpeng Tie, Haiyan Tang, Jiaquan Zhang)....Pages 93-103
Thermodynamic Study on Substitution of CO2 for Ar or O2 in AOD Smelting Process (Rongyue Wang, Zhangfu Yuan, Xiangtao Yu)....Pages 105-111
Front Matter ....Pages 113-113
Recent Progress on Metal Oxide Semiconductor Thin Film Transistor Application via Atomic Layer Deposition Method (Jiazhen Sheng, Jung-Hoon Lee, Tae-Hyun Hong, Wan-Ho Choi, Jin-Seong Park)....Pages 115-120
Adsorption of Fluoride Gases in Aluminum Production by Using of Nanotechnology (Mohsen Ameri Siahooei, Kambiz Bordbari)....Pages 121-136
Experimental Study on Competitive Adsorption of SF6 Decomposed Components on Nitrogen-Doped TiO2 Nanotubes Sensor (Jun Zhang, Xiaoxing Zhang, Hao Cui, GuoZhi Zhang)....Pages 137-142
Fabrication of Hardystonite Nano-bioceramic Coating on 306L Stainless Steel Substrate Using Electrophoretic Method and Evaluation of Its Corrosion Resistance to Improve Medical Performance (Iman Bagherpour)....Pages 143-154
Fabrication of Monodispersed Needle-Sized Hollow Core Polystyrene Microspheres (Stanley O. Omorogbe, Esther U. Ikhuoria, Hilary I. Ifijen, Aline Simo, Aireguamen Aigbodion, Malik Maaza)....Pages 155-164
Hydrangea-Like VS4 Microspheres: A Novel Structure Material for High-Performance Electrochemical Capacitor Electrode (Zheng-Wu Peng, Kai-Feng Jun, Hong-Yi Li, Bing Xie)....Pages 165-172
Preparation and Properties of Novel Graphene Composites (Wanlong Zhang, Haibin Zuo, Jingsong Wang, Yingli Liu, Yajie Wang)....Pages 173-183
Synthesis and Characterization of Silver Nanoparticles Using Simple Polyol Method (M. Tarek, A. M. El-Aziz)....Pages 185-194
Front Matter ....Pages 195-195
Differentiating Defect Types in LENSTM Metal AM via In Situ Pyrometer Process Monitoring (Tom Stockman, Caleb Horan, Cameron Knapp, Kevin Henderson, Brian Patterson, John Carpenter et al.)....Pages 197-204
Laser-Additive Repair of Cast Ni–Al–Bronze Components (Xinjin Cao, Priti Wanjara, Javad Gholipour, Yueping Wang)....Pages 205-216
Comparative Austempering Response Between Weld Metals of ADI Weldments With and Without Cerium Addition (Tapan Kumar Pal, Tapan Sarkar)....Pages 217-237
Effects of Beam Oscillation on Porosity and Intermetallic Compounds Formation of Electron Beam Welded DP600 Steel to Al-5754 Alloy Joints (Soumitra Kumar Dinda, Prakash Srirangam, Gour Gopal Roy)....Pages 239-249
Effects of Ultrasonic Micro-forging on 304 Stainless Steel Fabricated by WAAM (Laibo Sun, Fengchun Jiang, Ding Yuan, Xiaojing Sun, Yan Su, Chunhuan Guo)....Pages 251-258
Interface Microstructural Characterization of Titanium to Stainless Steel Dissimilar Friction Welds (Muralimohan Cheepu, V. Muthupandi, Woo Seong Che)....Pages 259-268
Mechanical Property Characterization of Single Scan Laser Tracks of Nickel Superalloy 625 by Nanoindentation (Jordan S. Weaver, Meir Kreitman, Jarred C. Heigel, M. Alkan Donmez)....Pages 269-278
Metallurgical Characteristics of Laser Peened 17-4 PH SS Processed by LENS Technique (I. Mathoho, E. T. Akinlabi, N. Arthur, M. Tlotleng, B. Masina)....Pages 279-285
Front Matter ....Pages 287-287
Prototyping of a Laboratory-Scale Cyclone Separator for Biofuel Production from Biomass Feedstocks Using a Fused Deposition Modeling Printer (Samuel Hansen, Amin Mirkouei)....Pages 289-297
Front Matter ....Pages 299-299
Phase-Field Modeling of Microstructure Evolution of Binary and Multicomponent Alloys During Selective Laser Melting (SLM) Process (Ali Ramazani, Julia Kundin, Christian Haase, Ulrich Prahl)....Pages 301-309
Phase-Field Simulation of Microstructure Evolution in Direct Metal Laser Sintered AlSi10Mg (Hossein Azizi, Nikolas Provatas, Mohsen Mohammadi)....Pages 311-318
Laser Interaction with Surface in Powder Bed Melting Process and Its Impact on Temperature Profile, Bead and Melt Pool Geometry (Leila Ladani, Faiyaz Ahsan)....Pages 319-329
Evolution of a Gradient Microstructure in Direct Metal Laser Sintered AlSi10Mg (Amir Hadadzadeh, Babak Shalchi Amirkhiz, Brian Langelier, Jian Li, Mohsen Mohammadi)....Pages 331-338
Finite Element Analysis of Particle Pushing During Selective Laser Melting of AlSi10Mg/AlN Composites (Marjan Nezafati, Ali Bakhshinejad, Benjamin Church, Pradeep Rohatgi)....Pages 339-346
Numerical Simulation on the Single-Crystal Grain Structure of GH4169 Superalloy Steel in the Spiral Grain Selector Using Procast Software (Zheng Chen, Lan’xin Geng, Yu Yao, Yi Cheng, Jieyu Zhang)....Pages 347-353
Powder Packing Density and Its Impact on SLM-Based Additive Manufacturing (Taher Abu-Lebdeh, Ransford Damptey, Vincent Lamberti, Sameer Hamoush)....Pages 355-367
Front Matter ....Pages 369-369
About a Digital Twin for the Fatigue Approach of Additively Manufactured Components (Rainer Wagener, Matilde Scurria, Thilo Bein)....Pages 371-382
Effect of the Surface Finish on the Cyclic Behavior of Additively Manufactured AlSi10Mg (Matilde Scurria, Benjamin Möller, Rainer Wagener, Tobias Melz)....Pages 383-394
Effect of Heat Treatments on Fatigue Properties of Ti–6Al–4V and 316L Produced by Laser Powder Bed Fusion in As-Built Surface Condition (Antonio Cutolo, Chola Elangeswaran, Charlotte de Formanoir, Gokula Krishna Muralidharan, Brecht Van Hooreweder)....Pages 395-405
Fracture Toughness and Fatigue Strength of Selective Laser Melted Aluminium–Silicon: An Overview (Leonhard Hitzler, Enes Sert, Markus Merkel, Andreas Öchsner, Ewald Werner)....Pages 407-412
The Effect of Heat Treatment and Alloying of Ni–Ti Alloy with Copper on Improving Its Fatigue Life (Wisam Abu Jadayil, Duaa Serhan)....Pages 413-420
Effect of Adding Yttrium on the Inclusion Modification and Impact Toughness of E36 Shipbuilding Steel (Xiaojun Xi, Maolin Ye, Shufeng Yang, Jingshe Li)....Pages 421-430
Front Matter ....Pages 431-431
Influence of Nitrogen on Microstructure, Mechanical Properties and Martensitic Phase Transformation of Co–26Cr–5Mo–5W Alloys by Selective Laser Melting (Bo Wang, Xinglong An, Fei Liu, Min Song, Song Ni, Shaojun Liu)....Pages 433-442
The Morphology, Crystallography, and Chemistry of Phases in Wire-Arc Additively Manufactured Nickel Aluminum Bronze (Chalasani Dharmendra, Amir Hadadzadeh, Babak Shalchi Amirkhiz, Mohsen Mohammadi)....Pages 443-453
Microstructure Evolution in Direct Metal Laser Sintered Corrax Maraging Stainless Steel (Amir Hadadzadeh, Babak Shalchi Amirkhiz, Jian Li, Mohsen Mohammadi)....Pages 455-462
The Microtexture and Tensile Properties of Continuous-Wave and Quasi-Continuous-Wave Laser Powder-Deposited Inconel 718 (Zhaoyang Liu, Qiang Zhu, Lijun Song)....Pages 463-471
Front Matter ....Pages 473-473
Alloy Design for Biomedical Applications in Additive Manufacturing (K.-P. Hoyer, M. Schaper)....Pages 475-484
Surface Inoculation of Aluminium Powders for Additive Manufacturing Guided by Differential Fast Scanning Calorimetry (Lennart Tasche, Kay-Peter Hoyer, Evgeny Zhuravlev, Guido Grundmeier, Mirko Schaper, Olaf Keßler)....Pages 485-493
Mechanical Behavior and Microstructure of Porous Ti Using TiC as Reinforcement (Shiyuan Liu, Jian Wang, Tengfei Lu, Guibao Qiu, Hao Cui)....Pages 495-501
Processing of Haynes® 282® Alloy by Laser Powder Bed Fusion Technology (Robert Otto, Vegard Brøtan, Amin S. Azar, Olav Åsebø)....Pages 503-510
Front Matter ....Pages 511-511
Tensile Deformation Behavior of 1 GPa-Grade TRIP-Aided Multi-microstructure Steels Studied by In Situ Neutron Diffraction (Noriyuki Tsuchida, Takaaki Tanaka, Yuki Toji)....Pages 513-518
Development of Advanced High-Strength Steels for Automobile Applications (Francys Barrado, Tihe Zhou, David Overby, Peter Badgley, Chris Martin-Root, Sarah Zhang et al.)....Pages 519-527
Effect of Carbon Content on Strengthening Behavior with Grain Refinement on Lath Martensite Structure (Hiroyuki Kawata, Yoshiaki Honda, Kengo Takeda)....Pages 529-535
Assessment of the Strengthening Mechanisms Operating in Microalloyed Steels During Cyclic Deformation Using High-Resolution Electron Backscatter Diffraction (Paulina Lisiecka-Graca, Krzysztof Muszka, Janusz Majta)....Pages 537-547
Effect of Niobium on Microstructure and Mechanical Properties of Nb–Ti Microalloyed Carbide-Free Bainitic Steels (Xi Chen, Fuming Wang, Changrong Li, Shuai Liu)....Pages 549-560
Effect of Inclusions Modified by Y-Based Rare Earth on the Corrosion Behavior of EH36 Shipbuilding Steel (Maolin Ye, Xiaojun Xi, Libin Zhu, Shufeng Yang, Jingshe Li)....Pages 561-569
Microstructure and Mechanical Properties of Intercritical Annealed Multiphase Ultrahigh Strength Steel (Huasai Liu, Xiangyu Li, Chunqian Xie, Yun Han)....Pages 571-577
The Effect of Ni and Cu Addition on Mechanical Behavior of Thermomechanically Controlled Processed HSLA X100 Steels (A. R. Hosseini Far, S. H. Mousavi Anijdan, M. Abbasi)....Pages 579-590
Front Matter ....Pages 591-591
Optimization of Magnetocaloric Properties of Ball-Milled La(Fe,Co,Si)\(_{13}\)(H,C)\(_y\) (V. Paul-Boncour, K. Nakouri, L. Bessais)....Pages 593-598
Production of High-Resistivity Electrical Steel Alloys by Substitution of Si with Al and Cr (Brhayan Stiven Puentes Rodriguez, David Brice, James B. Mann, Srinivasan Chandrasekar, Kevin Trumble)....Pages 599-606
Nanocrystalline Multifunctional Pr–Co Compounds (W. Bouzidi, T. Bartoli, A. Michalowicz, J. Moscovici, N. Mliki, L. Bessais)....Pages 607-615
Front Matter ....Pages 617-617
A Study on Electrical Conductivity of Micro Friction Stir-Welded Dissimilar Sheets for Hybrid Electric Vehicles (HEVs) (Omkar Mypati, Surjya Kanta Pal, Prakash Srirangam)....Pages 619-627
Micro-structure and Properties of Cu–0.3 wt%Ag Alloy Ultra-Fine Wires (Shu-sen Wang, Yuan-wang Zhang, Da-wei Yao)....Pages 629-635
Length Scale of the Cellular Microstructure Tailoring Tensile Properties of Zn–20 wt%Sn–2 wt%Cu Solder Alloy (Cesar Bertolin dos Santos Mangualde, Rodrigo Valenzuela Reyes, José Eduardo Spinelli)....Pages 637-644
Effect of Ag on the Mechanical Properties of Bi–Ag Solder Alloys by the Single-Lap Shear Test Method (Nima Ghamarian, M. A. Azmah Hanim, M. Nahavandi, Ali Ourdjini, Zulkarnain Zainal, H. N. Lim)....Pages 645-653
Front Matter ....Pages 655-655
Parametrically Homogenized Continuum Damage Mechanics (PHCDM) Models for Composites from Micromechanical Analysis (Xiaofan Zhang, Zhiye Li, Daniel J. O’Brien, Somnath Ghosh)....Pages 657-665
Effect of Multi-gating System on Solidification of Molten Metals in Spur Gear Casting: A Simulation Approach (Enesi Y. Salawu, Emuowhochere Oghenevwegba, Oluseyi O. Ajayi, A. O. Inegbenebor, E. T. Akinlabi, S. T. Akinlabi)....Pages 667-677
Front Matter ....Pages 679-679
Corrosion Study of Boron Nitride Nanosheets Deposited on Copper Metal by Electrophoretic Deposition (Mohsin Ali Raza, Amer Nadeem, Muhammad Tasaduq Ilyas)....Pages 681-685
Effects of Process Parameters on the Zirconia Coating Prepared by Sol-Gel and Electrodeposition Process (Jian Dong, Yanhui Sun, Bingsheng Dou, Feiyu He, Hongtao Huang, Jianping Zhen)....Pages 687-696
The Study of Slurry Erosion Wear Behaviour of Coal Bottom Ash Slurry Handling Pipeline (Satish R. More, Sudeep P. Ingole, Dhananjay V. Bhatt, Jyoti V. Menghani)....Pages 697-710
Wear Characterization of Cemented Carbide Multipoint Cutting Tool Machining AISI 4140 at High Cutting Speed: Criteria for Materials Selection (Federico Simone Gobber, Elisa Fracchia, Mario Rosso)....Pages 711-718
Dry Sheet Metal Forming Through Selective Oxidized Tool Surfaces (Bernd-Arno Behrens, Deniz Yilkiran, Simon Schöler, Sven Hübner, Kai Möhwald, Fahrettin Özkaya)....Pages 719-731
Effect of Process Parameters on Surface Properties of Laser-Hardened Cast Iron (S. V. Wagh, Sudeep Ingole, D. V. Bhatt, J. V. Menghani, M. J. Rathod)....Pages 733-743
On Improvement in Surface Integrity of µ-EDMed Ti–6Al–4V Alloy by µ-ECM Process ( Ramver, Akshay Dvivedi, Pradeep Kumar)....Pages 745-753
Corrosion and Wear Resistance of PTFE-Al2O3 Coatings Deposited on Aluminum Alloy by a Microblasting Process (A. M. Oladoye, J. G. Carton, A. Baroutaji, M. Obeidi, J. Stokes, B. Twomey et al.)....Pages 755-762
Front Matter ....Pages 763-763
Numerical Simulation of Ti6–Al4–V Alloy Diffusion Bonding Process Based on Molecular Dynamics (Xiaogang Liu, Yongji Zuo, Haiding Guo)....Pages 765-777
Front Matter ....Pages 779-779
Custom Pyrolytic Graphite–Steel Thermocouple for High-Temperature Measurements (Abdul-Sommed Hadi, Bryce E. Hill)....Pages 781-789
Front Matter ....Pages 791-791
3D Contact and Strain in Alveolar Bone Under Tooth/Implant Loading (Yuxiao Zhou, Chujie Gong, Mehran Hossaini-Zadeh, Jing Du)....Pages 793-798
Shear-Punch Testing of Human Cranial Bone and Surrogate Materials (A. D. Brown, C. A. Gunnarsson, K. A. Rafaels, S. Alexander, T. A. Plaisted, T. Weerasooriya)....Pages 799-808
Investigation of Biodegradable Zn–Li–Cu Alloys for Orthopaedic and Cardiovascular Applications (Jacob Young, Ramana G. Reddy)....Pages 809-818
Low-Temperature Air Plasma Modification of Electrospun Soft Materials and Bio-interfaces (Bernabe S. Tucker, Ranu Surolia, Paul A. Baker, Yogesh Vohra, Veena Antony, Vinoy Thomas)....Pages 819-826
Accumulation of Biofilm on Ti–6Al–4V Alloy Fabricated Using Additive Layer Manufacturing (Mari Koike, Tetsuro Horie, Richard J. Mitchell, Toru Okabe)....Pages 827-836
Copper Recovery from Printed Circuit Boards from Smartphones Through Bioleaching (Lidiane Maria de Andrade, Carlos Gonzalo Alvarez Rosario, Mariana Alves de Carvalho, Denise Crocce Romano Espinosa, Jorge Alberto Soares Tenório)....Pages 837-844
Dependence of the Ferrovanadium Power as Additive on Mechanical Property in Porous Ti (Guibao Qiu, Jian Wang, Shiyuan Liu, Chenguang Bai, Yilong Liao)....Pages 845-854
Effect of Compaction Pressure on Porosity and Mechanical Properties of Porous Titanium as Bone Substitute Materials (Qingjuan Li, Guibao Qiu, Shiyuan Liu, Tengfei Lu)....Pages 855-864
The Effect of Milling Time on Structural, Friction and Wear Behavior of Hot Isostatically Pressed Ti–Ni Alloys for Orthopedic Applications (Mamoun Fellah, Naouel Hezil, Mohammed Abdul Samad, Mohamed Zine Touhami, Alex Montagne, Alain Iost et al.)....Pages 865-875
Front Matter ....Pages 877-877
Perturbation Analysis of Amorphous Alloy Formation (Rahul Basu)....Pages 879-885
Shockwave Consolidation to Create Bulk Metallic Glass (David Nemir, Jan Beck, Lawrence Murr, Yirong Lin, Luis Chavez)....Pages 887-897
Front Matter ....Pages 899-899
Characterization of the Irradiation Effects in Nuclear Graphite (J. David Arregui-Mena, Philip D. Edmondson, Robert N. Worth, Cristian Contescu, Timothy D. Burchell, Yutai Katoh)....Pages 901-906
Irradiation Effects on Reactor Concrete Structures (J. David Arregui-Mena, Alain B. Giorla, G. E. Jellison, Elena Tajuelo-Rodriguez, Christa E. Torrence, Masaki Kawai et al.)....Pages 907-912
Front Matter ....Pages 913-913
Electrochemical Mechanism and Preparation of Cr–Low-Carbon Steel Composite in a NaCl–KCl–NaF–Cr2O3 Molten Salt (Shixian Zhang, Yungang Li, Cong Wang, Xiaoping Zhao)....Pages 915-926
Diamond-Like Carbon Coating for Drill Collars: Test Experiences (Nausha Asrar, Jeffrey Ham)....Pages 927-937
Inhibition Effect of Essential Oil Extracts on the Corrosion Inhibition of Mild Steel in Chloride–Sulphate Media (Roland Tolulope Loto, Richard Leramo, Babatunde Oyebade)....Pages 939-948
Corrosion Properties of Steel Sheet with Zinc-Base Alloy Coatings (Guangrui Jiang, Guanghui Liu, Ting Shang, Wanling Qiu)....Pages 949-957
Effect of Heat Treatment on the Localized Corrosion Resistance of S32101 Duplex Stainless Steel in Chloride/Sulphate Media (Roland Tolulope Loto, Cleophas Akintoye Loto, Akanji Olaitan, Olufunmilola Joseph)....Pages 959-966
Study of Mechanisms of Cobalt Electrodeposition by Means of Potentiodynamic Polarization Curves (M. Ohba, T. Scarazzato, D. C. R. Espinosa, J. A. S. Tenório, Z. Panossian)....Pages 967-976
Front Matter ....Pages 977-977
Artificial Intelligent and Simulation Nanostructure of Ceramic (Habibollah Aminirastabi, Fatemeh Karimidehcheshmeh, Gouli Ji)....Pages 979-993
Front Matter ....Pages 995-995
Kinetics Calculation and Analysis of AlN Precipitation in ML40Cr Steel Austenite (Ziyi Liu, Yanping Bao, Min Wang)....Pages 997-1005
Study of Dendrite Growth Under Forced Convection in Superalloy Solidification by Multiphase-Field Coupled Lattice Boltzmann Method (Cong Yang, Qingyan Xu, Baicheng Liu)....Pages 1007-1017
Modeling of Volume Diffusion-Controlled Phase Transformations in Multiphase Multicomponent Alloy Systems by Minimization of Gibbs Energy (Anders Salwén)....Pages 1019-1025
A New Method for Calculation of Vapor–Liquid Equilibrium (VLE) of Au–Cu Alloy System (Lingxin Kong, Jingbao Gao, Junjie Xu, Baoqiang Xu, Bin Yang, Yifu Li)....Pages 1027-1035
Ab Initio Study on the Oxidation Mechanism of Millerite (Xiaolu Xiong, Xionggang Lu, Guangshi Li, Hongwei Cheng, Qian Xu, Shenggang Li)....Pages 1037-1044
Kinetic Model of Silica Dissolution in CaO–SiO2–MgO–Al2O3 Slag System (Haifei An, Jie Li, Aimin Yang, Weixing Liu, Can Tian)....Pages 1045-1053
Front Matter ....Pages 1055-1055
Origin of the Significant Impact of Ta on the Creep Resistance of FeCrNi Alloys (D. Magne, X. Sauvage, M. Couvrat)....Pages 1057-1066
Stress Analysis and Structure Optimization of W-Shaped Radiant Tube in Continuous Annealing Furnace (Yang Long Li, Shun Ming Liu, Da Wei Hou, Wei Guo, Hui Wang, Meng Yu)....Pages 1067-1075
Front Matter ....Pages 1077-1077
Case Studies of Continuous Improvement Projects in the Metals Industry (Cynthia Belt)....Pages 1079-1086
Front Matter ....Pages 1087-1087
Fracture Mechanics-Based Study of Stress Corrosion Cracking of SS304 Dry Storage Canister for Spent Nuclear Fuel (Leonardi Tjayadi, Nilesh Kumar, Korukonda L. Murty)....Pages 1089-1097
Similar and Dissimilar Metal Weld Failures in Hydrocracking Service at a Refinery (Sudhakar Mahajanam, Cesar Espinoza, Yenny Cubides)....Pages 1099-1109
Influence of Tempering Treatment on Precipitation Behavior, Microstructure, Dislocation Density and Hydrogen-Induced Ductility Loss in High-Vanadium Hot-Rolled X80 Pipeline Steel (Longfei Li, Bo Song, Zeyun Cai, Zhen Liu, Xiaokang Cui)....Pages 1111-1122
Front Matter ....Pages 1123-1123
Initiation and Early Growth of Fatigue Cracks (Jaroslav Polák)....Pages 1125-1135
Front Matter ....Pages 1137-1137
Influence of CaO/SiO2/Al2O3 Ratio on the Melting Behaviour of SynCon Slags (Dominik Hofer, Stefan Luidold, Tobias Beckmann, Frank Schulenburg)....Pages 1139-1148
Freeze Lining Refractories in Non-ferrous TSL Smelting Systems (Stanko Nikolic, Ben Hogg, Paul Voigt)....Pages 1149-1159
Freeze-Lining Formation in Submerged Arc Furnaces Producing Ferrochrome Alloy in South Africa (Joalet Dalene Steenkamp, Quinn Gareth Reynolds, Markus Wouter Erwee, Stefan Swanepoel)....Pages 1161-1180
Designing Furnace Lining/Cooling Systems to Operate with a Competent Freeze Lining (Hugo Joubert, Isobel Mc Dougall)....Pages 1181-1195
Front Matter ....Pages 1197-1197
Corrosion Resistance of Hot Dipping Al–Zn–Si and Zn–Al–Mg–Si Alloy Coating (Hui Li, Jinglong Liang, Dongbin Wang, Yungang Li)....Pages 1199-1206
Performance of Low-Cost 3D Printed Pylon in Lower Limb Prosthetic Device (Fariborz Tavangarian, Camila Proano, Caleb Zolko)....Pages 1207-1215
Sequential Leaching Characteristics of Chromium in AOD Slag-Based Cementitious Materials (Ya-Jun Wang, Jun-Guo Li, Ya-Nan Zeng, Xiao-Yu Li)....Pages 1217-1225
Study on the Reaction Behavior of Hydrochloric-Acid-Containing Titanium Blast Furnace Slag (Jinglong Liang, Hui Li, Jing Wang, Dongbin Wang, Ramana G. Reddy, Yu Yang)....Pages 1227-1235
Thermodynamic and Kinetic Analysis of Inhomogeneous Distribution of Solute on Precipitations in as Cast Nb–V–Ti Microalloyed Steel (Ya-Nan Zeng, Jun-Guo Li, Ya-Jun Wang)....Pages 1237-1247
Front Matter ....Pages 1249-1249
Radiation and Corrosion Resistances of 316LN Austenitic Stainless Steel by Rotationally Accelerated Shot Peening (Bin Yang, Xudong Chen, Yuntian Zhu, Yusheng Li)....Pages 1251-1260
Front Matter ....Pages 1261-1261
A Novel Dual-Phase Gradient Material of High-Entropy Alloy Prepared by Spark Plasma Sintering (Wei Zhang, Mingyang Zhang, Fangzhou Liu, Yingbo Peng, Songhao Hu, Yong Liu)....Pages 1263-1270
Molecular Dynamics Simulations on the Mechanical Behavior of AlCoCrCu0.5FeNi High-Entropy Alloy Nanopillars (Wei Li, Jing Tang, Qingyuan Wang, Haidong Fan)....Pages 1271-1280
Production of AlCoCrFeNiME-Based High-Entropy Alloys via Self-Propagating High-Temperature Synthesis (Murat Alkan, Esra Dokumaci, Berkay Türkoglu, Aslihan Kara, Büsra Aksu, Dilan Ugurluer)....Pages 1281-1287
Front Matter ....Pages 1289-1289
ICME Applied in the Undergraduate Capstone Senior Design Sequence (Paul Sanders)....Pages 1291-1301
Front Matter ....Pages 1303-1303
Effect of a Vertical Twin Boundary on the Mechanical Property of Bicrystalline Copper Micropillars (DeAn Wei, Haidong Fan, Jing Tang, Xu Zhang)....Pages 1305-1310
Front Matter ....Pages 1311-1311
Effect of Ni on the Corrosion Behavior of Haynes 230 Alloy in MgCl2-KCl Salt (Yuxiang Peng, Ramana G. Reddy)....Pages 1313-1321
Front Matter ....Pages 1323-1323
The Study of Mechanical Behaviour of Materials for the Nuclear Reactor Components in SUSEN Hot Cells (Mariia Zimina, Petr Švrčula, Pavel Zháňal, Ondřej Libera, Stefan Zaunschirm, Ondřej Srba)....Pages 1325-1333
Investigation of Radiation Temperature and Straining Temperature Effects on the Screw Dislocation Mobility Evolution in Irradiated Ferritic Grains Using 3D Dislocation Dynamics (Yang Li, Christian Robertson, Xianfeng Ma, Biao Wang)....Pages 1335-1344
Front Matter ....Pages 1345-1345
Mechanical Properties of Amorphous Silicon Nanoparticles (D. Kilymis, C. Gerard, L. Pizzagalli)....Pages 1347-1354
Front Matter ....Pages 1355-1355
Molecular Dynamics Simulation of the Structure and Transport Properties of xKF–yNaF–zAlF3 (Jie Li, Hui Guo, Hongliang Zhang, Ru Cai Li, Qiyu Wang, Jingkun Wang et al.)....Pages 1357-1369
Front Matter ....Pages 1371-1371
Microstructure Evolution and Physics Properties of Low Silver Copper Alloy Wires During In Situ Composite Preparation (Yuan-wang Zhang, Shu-sen Wang, Da-wei Yao)....Pages 1373-1380
Front Matter ....Pages 1381-1381
Heat Treatment Strategies to Improve the Quasi-Static and Dynamic Performance of Alpha + Beta Titanium Alloys (Alireza Fadavi Boostani, Shiraz Mujahid, Andrew L. Oppedal, Cory Krivanec, Wilburn R. Whittington, Paul G. Allison et al.)....Pages 1383-1387
Effect of Silicon Content on the Dilatometric Behavior of a Medium-Carbon Steel (A. I. Gallegos Pérez, O. Vázquez Gómez, J. J. López Soria, H. J. Vergara Hernández, E. López Martínez)....Pages 1389-1400
Phase-Field Simulation of Intermetallic Phase Precipitation in a High-Al Alloyed Lightweight High-Strength Steel (Carsten Drouven, Bowen Zou, Wenwen Song, Wolfgang Bleck)....Pages 1401-1409
Shape Memory Behavior of Ni49.5Ti50.5 Processing-Induced Strain Glass Alloys (Robert W. Wheeler, Jesse Smith, Nathan A. Ley, Anit Giri, Marcus L. Young)....Pages 1411-1420
Precipitation Hardening of Supersaturated Al–Sc–Zr Produced via Melt-Spinning (Yang Yang, Paul Sanders)....Pages 1421-1426
Effect of Sm Content and Solidification Rate on Microstructure of SmFe Alloy (Kun Liu, Shuhuan Wang, Yunli Fen, Chunyan Song, Guolong Ni, Kaixuan Zhang)....Pages 1427-1436
Evolution of Dendritic Morphology Under HPMO Treatment (Hui-cheng Li, Yu-xiang Liu, Zhen Liu, Qi-jie Zhai)....Pages 1437-1445
In Situ Observation of Melting and δ ↔ γ Phase Transformation in Duplex Stainless Steel (Yang Liu, Yan-hui Sun)....Pages 1447-1457
Microstructural Evolution of a Transformation in Which There Is an Exclusion Zone Around Each Nucleus (Paulo R. Rios, Harison S. Ventura, André L. M. Alves, Weslley L. S. Assis, Elena Villa)....Pages 1459-1469
Thermodynamic Properties of Si–B Alloys Determined by Solid-State Heterogeneous Phase Equilibrium (Muhammad A. Imam, Ramana G. Reddy)....Pages 1471-1479
Front Matter ....Pages 1481-1481
Microstructure Evolution and Mechanical Properties of Medical Material Mg–3Zn Alloy Prepared by Semi-solid Powder Injection Moulding (Xia Luo, Chao Fang, Zhou Fan, Bensheng Huang, Jun Yang)....Pages 1483-1497
Inhomogeneity of Strain in Metal Particulates Produced by Modulation-Assisted Machining (Indrani Biswas, James B. Mann, Srinivasan Chandrasekar, Kevin Trumble)....Pages 1499-1506
Numerical Simulation and Validation of Gas and Molten Metal Flows in Close-Coupled Gas Atomization (F. Hernandez, T. Riedemann, J. Tiarks, B. Kong, J. D. Regele, T. Ward et al.)....Pages 1507-1519
Density Separation of Mixed Carbide Colloids via Standing Wave Physics (Trenin K. Bayless, Jerome P. Downey, Grant C. Wallace, Mark D’Aberle)....Pages 1521-1531
The Influence of Mechanical Activation on the Synthesis of Ca2MgSi2O7 (Fariborz Tavangarian, Caleb Zolko)....Pages 1533-1541
Front Matter ....Pages 1543-1543
3D Printing of Polymer-Based Gasochromic, Thermochromic and Piezochromic Sensors (Patrick Dzisah, Airefetalo Sadoh, Nuggehalli M. Ravindra)....Pages 1545-1561
3D Printing of Pharmaceuticals and Transdermal Drug Delivery––An Overview (David Bird, Emel Eker, Nuggehalli M. Ravindra)....Pages 1563-1573
Formulation of Curable Resins Utilized in Stereolithography (David Bird, Elbert Caravaca, Joseph Laquidara, Keith Luhmann, Nuggehalli M. Ravindra)....Pages 1575-1587
MARS––Magnetic Augmented Rotation System (Vishwas Danthi Shivaram, Roulei Liu, Navjot Panchhi, Laila Alqarni, Rayan Daroowalla, Shuang Du et al.)....Pages 1589-1600
Front Matter ....Pages 1601-1601
Friction Conditions on Deep-Drawing Tool Radii When Using Volatile Media as Lubrication Substitute (Gerd Reichardt, Mathias Liewald)....Pages 1603-1613
Investigation of Friction and Adhesion Behavior of Textured Workpieces and Coated Tools Under Dry Tribological Contact (Rafael Hild, Robby Mannens, Daniel Trauth, Patrick Mattfeld, Thomas Bergs, Dennis C. Hoffmann et al.)....Pages 1615-1628
Effects of Emissivity on Combustion Behavior of Energetic Materials (Elbert Caravaca, David Bird, Henry Grau, Viral Panchal, Nuggehalli M. Ravindra)....Pages 1629-1641
Self-healing in Materials: An Overview (Samiha Hossain, Nuggehalli M. Ravindra)....Pages 1643-1661
Front Matter ....Pages 1663-1663
Revealing the Heterogeneous Nucleation and Growth Behaviour of Grains in Inoculated Aluminium Alloys During Solidification (Yijiang Xu, Daniele Casari, Ragnvald H. Mathiesen, Yanjun Li)....Pages 1665-1675
Influence of Microstructure Evolution During Twin-Roll Casting on the Properties of Magnesium Sheets (K. U. Kainer, G. Kurz, S. Pakulat, D. Letzig)....Pages 1677-1686
A History of the Global Light Metals Alliance (Jennifer Jackman, Kumar Sadayappan, Mark Easton)....Pages 1687-1696
Analysis of the High-Purity Aluminum Purification Process Using Zone-Refining Technique (Heli Wan, Baoqiang Xu, Jinyang Zhao, Bin Yang, Yongnian Dai)....Pages 1697-1706
Back Matter ....Pages 1707-1731

Citation preview

SUPPLEMENTAL PROCEEDINGS

The Minerals, Metals & Materials Series

The Minerals, Metals & Materials Society Editor

TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings

123

Editor The Minerals, Metals & Materials Society Pittsburgh, PA, USA

ISSN 2367-1181 ISSN 2367-1696 (electronic) The Minerals, Metals & Materials Series ISBN 978-3-030-05860-9 ISBN 978-3-030-05861-6 (eBook) https://doi.org/10.1007/978-3-030-05861-6 Library of Congress Control Number: 2018964042 © The Minerals, Metals & Materials Society 2019 This work is subject to copyright. All rights are reserved by the Publisher, whether the whole or part of the material is concerned, specifically the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfilms or in any other physical way, and transmission or information storage and retrieval, electronic adaptation, computer software, or by similar or dissimilar methodology now known or hereafter developed. The use of general descriptive names, registered names, trademarks, service marks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. The publisher, the authors and the editors are safe to assume that the advice and information in this book are believed to be true and accurate at the date of publication. Neither the publisher nor the authors or the editors give a warranty, express or implied, with respect to the material contained herein or for any errors or omissions that may have been made. The publisher remains neutral with regard to jurisdictional claims in published maps and institutional affiliations. This Springer imprint is published by the registered company Springer Nature Switzerland AG The registered company address is: Gewerbestrasse 11, 6330 Cham, Switzerland

Contents

Part I

2019 International Metallurgical Processes Workshop for Young Scholars (IMPROWYS 2019)

Ab Initio Molecular Dynamics Study on the Dissolution of Interfacial Iron Oxides in Hot Compressive Bonding Combined with Experiments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Honglin Zhang, Mingyue Sun, Bin Xu and Dianzhong Li Effect of MgO Content on the Properties of Magnesia Fluxed Pellets . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Yuzhu Zhang, Weixing Liu, Aimin Yang and Jie Li Numerical Simulation of Three-Phase Flow of Gas-Stirring Micro-phenomenon During Ladle Furnace Process . . . . . . . . . . . . . . Libin Zhu, Wei Liu, Shfueng Yang, Jingshe Li, Feng Wang and Xueliang Zhang The Effect of pH and Temperature During Carbonation Process on Spent Die Cleaning Solution from Aluminium Extrusion Industry . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Ahmed S. Aadli Improvement of Center Segregation in Continuously Cast Blooms by Convex Roll Soft Reduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . Liang Li, Xiao Zhao, Peng Lan, Zhanpeng Tie, Haiyan Tang and Jiaquan Zhang Effects of a Top-Down Flow on Gas–Solid Fluidization State in a Bubble Fluidized Bed . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Xu Han, Liangying Wen, Shengyun Shi, Wenhuan Jiang, Meihuan Liu and Feng Lu

3

17

29

39

51

63

v

vi

Contents

Development of Bio-treated Oil Palm Fiber Reinforced Kaolin Matrix Composites for Building Bricks Application . . . . . . . . . . . . . . . . . . . Muideen Adebayo Bodude, Olasunkanmi B. Adegbuyi and Ruth Nkiruka Nnaji Effect of Roll Surface Profile on Thermal-Mechanical Behavior of Continuously Cast Bloom in Soft Reduction Process . . . . . . . . . . . Liang Li, Xiao Zhao, Peng Lan, Zhanpeng Tie, Haiyan Tang and Jiaquan Zhang Thermodynamic Study on Substitution of CO2 for Ar or O2 in AOD Smelting Process . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Rongyue Wang, Zhangfu Yuan and Xiangtao Yu Part II

77

93

105

2019 Symposium on Functional Nanomaterials: Synthesis, Integration, and Application of Emerging Nanomaterials

Recent Progress on Metal Oxide Semiconductor Thin Film Transistor Application via Atomic Layer Deposition Method . . . . . . . . . . . . . . . Jiazhen Sheng, Jung-Hoon Lee, Tae-Hyun Hong, Wan-Ho Choi and Jin-Seong Park

115

Adsorption of Fluoride Gases in Aluminum Production by Using of Nanotechnology . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Mohsen Ameri Siahooei and Kambiz Bordbari

121

Experimental Study on Competitive Adsorption of SF6 Decomposed Components on Nitrogen-Doped TiO2 Nanotubes Sensor . . . . . . . . . . Jun Zhang, Xiaoxing Zhang, Hao Cui and GuoZhi Zhang

137

Fabrication of Hardystonite Nano-bioceramic Coating on 306L Stainless Steel Substrate Using Electrophoretic Method and Evaluation of Its Corrosion Resistance to Improve Medical Performance . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Iman Bagherpour Fabrication of Monodispersed Needle-Sized Hollow Core Polystyrene Microspheres . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Stanley O. Omorogbe, Esther U. Ikhuoria, Hilary I. Ifijen, Aline Simo, Aireguamen Aigbodion and Malik Maaza Hydrangea-Like VS4 Microspheres: A Novel Structure Material for High-Performance Electrochemical Capacitor Electrode . . . . . . . . . . Zheng-Wu Peng, Kai-Feng Jun, Hong-Yi Li and Bing Xie Preparation and Properties of Novel Graphene Composites . . . . . . . . Wanlong Zhang, Haibin Zuo, Jingsong Wang, Yingli Liu and Yajie Wang

143

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vii

Synthesis and Characterization of Silver Nanoparticles Using Simple Polyol Method . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . M. Tarek and A. M. El-Aziz Part III

185

Additive Manufacturing and Welding: Physical and Mechanical Metallurgy of Rapidly Solidified Metals

Differentiating Defect Types in LENSTM Metal AM via In Situ Pyrometer Process Monitoring . . . . . . . . . . . . . . . . . . . . . . . . . . . . Tom Stockman, Caleb Horan, Cameron Knapp, Kevin Henderson, Brian Patterson, John Carpenter and Judith Schneider Laser-Additive Repair of Cast Ni–Al–Bronze Components . . . . . . . . . Xinjin Cao, Priti Wanjara, Javad Gholipour and Yueping Wang Comparative Austempering Response Between Weld Metals of ADI Weldments With and Without Cerium Addition . . . . . . . . . . . . . . . . Tapan Kumar Pal and Tapan Sarkar Effects of Beam Oscillation on Porosity and Intermetallic Compounds Formation of Electron Beam Welded DP600 Steel to Al-5754 Alloy Joints . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Soumitra Kumar Dinda, Prakash Srirangam and Gour Gopal Roy Effects of Ultrasonic Micro-forging on 304 Stainless Steel Fabricated by WAAM . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Laibo Sun, Fengchun Jiang, Ding Yuan, Xiaojing Sun, Yan Su and Chunhuan Guo

197

205

217

239

251

Interface Microstructural Characterization of Titanium to Stainless Steel Dissimilar Friction Welds . . . . . . . . . . . . . . . . . . . . . . . . . . . . Muralimohan Cheepu, V. Muthupandi and Woo Seong Che

259

Mechanical Property Characterization of Single Scan Laser Tracks of Nickel Superalloy 625 by Nanoindentation . . . . . . . . . . . . . . . . . . Jordan S. Weaver, Meir Kreitman, Jarred C. Heigel and M. Alkan Donmez

269

Metallurgical Characteristics of Laser Peened 17-4 PH SS Processed by LENS Technique . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . I. Mathoho, E. T. Akinlabi, N. Arthur, M. Tlotleng and B. Masina

279

Part IV

Additive Manufacturing for Energy Applications

Prototyping of a Laboratory-Scale Cyclone Separator for Biofuel Production from Biomass Feedstocks Using a Fused Deposition Modeling Printer . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Samuel Hansen and Amin Mirkouei

289

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Part V

Contents

Additive Manufacturing of Metals: Applications of Solidification Fundamentals

Phase-Field Modeling of Microstructure Evolution of Binary and Multicomponent Alloys During Selective Laser Melting (SLM) Process . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Ali Ramazani, Julia Kundin, Christian Haase and Ulrich Prahl

301

Phase-Field Simulation of Microstructure Evolution in Direct Metal Laser Sintered AlSi10Mg . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Hossein Azizi, Nikolas Provatas and Mohsen Mohammadi

311

Laser Interaction with Surface in Powder Bed Melting Process and Its Impact on Temperature Profile, Bead and Melt Pool Geometry . . . . . Leila Ladani and Faiyaz Ahsan

319

Evolution of a Gradient Microstructure in Direct Metal Laser Sintered AlSi10Mg . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Amir Hadadzadeh, Babak Shalchi Amirkhiz, Brian Langelier, Jian Li and Mohsen Mohammadi Finite Element Analysis of Particle Pushing During Selective Laser Melting of AlSi10Mg/AlN Composites . . . . . . . . . . . . . . . . . . . . . . . Marjan Nezafati, Ali Bakhshinejad, Benjamin Church and Pradeep Rohatgi Numerical Simulation on the Single-Crystal Grain Structure of GH4169 Superalloy Steel in the Spiral Grain Selector Using Procast Software . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Zheng Chen, Lan’xin Geng, Yu Yao, Yi Cheng and Jieyu Zhang Powder Packing Density and Its Impact on SLM-Based Additive Manufacturing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Taher Abu-Lebdeh, Ransford Damptey, Vincent Lamberti and Sameer Hamoush Part VI

331

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347

355

Additive Manufacturing of Metals: Fatigue and Fracture III

About a Digital Twin for the Fatigue Approach of Additively Manufactured Components . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Rainer Wagener, Matilde Scurria and Thilo Bein

371

Effect of the Surface Finish on the Cyclic Behavior of Additively Manufactured AlSi10Mg . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Matilde Scurria, Benjamin Möller, Rainer Wagener and Tobias Melz

383

Contents

ix

Effect of Heat Treatments on Fatigue Properties of Ti–6Al–4V and 316L Produced by Laser Powder Bed Fusion in As-Built Surface Condition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Antonio Cutolo, Chola Elangeswaran, Charlotte de Formanoir, Gokula Krishna Muralidharan and Brecht Van Hooreweder Fracture Toughness and Fatigue Strength of Selective Laser Melted Aluminium–Silicon: An Overview . . . . . . . . . . . . . . . . . . . . . . . . . . Leonhard Hitzler, Enes Sert, Markus Merkel, Andreas Öchsner and Ewald Werner

395

407

The Effect of Heat Treatment and Alloying of Ni–Ti Alloy with Copper on Improving Its Fatigue Life . . . . . . . . . . . . . . . . . . . . . . . Wisam Abu Jadayil and Duaa Serhan

413

Effect of Adding Yttrium on the Inclusion Modification and Impact Toughness of E36 Shipbuilding Steel . . . . . . . . . . . . . . . . . . . . . . . . Xiaojun Xi, Maolin Ye, Shufeng Yang and Jingshe Li

421

Part VII

Additive Manufacturing of Metals: Microstructural Evolution and Phase Transformations

Influence of Nitrogen on Microstructure, Mechanical Properties and Martensitic Phase Transformation of Co–26Cr–5Mo–5W Alloys by Selective Laser Melting . . . . . . . . . . . . . . . . . . . . . . . . . . Bo Wang, Xinglong An, Fei Liu, Min Song, Song Ni and Shaojun Liu The Morphology, Crystallography, and Chemistry of Phases in Wire-Arc Additively Manufactured Nickel Aluminum Bronze . . . . Chalasani Dharmendra, Amir Hadadzadeh, Babak Shalchi Amirkhiz and Mohsen Mohammadi Microstructure Evolution in Direct Metal Laser Sintered Corrax Maraging Stainless Steel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Amir Hadadzadeh, Babak Shalchi Amirkhiz, Jian Li and Mohsen Mohammadi The Microtexture and Tensile Properties of Continuous-Wave and Quasi-Continuous-Wave Laser Powder-Deposited Inconel 718 . . . . . . Zhaoyang Liu, Qiang Zhu and Lijun Song Part VIII

433

443

455

463

Additive Manufacturing: Materials Design and Alloy Development

Alloy Design for Biomedical Applications in Additive Manufacturing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . K.-P. Hoyer and M. Schaper

475

x

Contents

Surface Inoculation of Aluminium Powders for Additive Manufacturing Guided by Differential Fast Scanning Calorimetry . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Lennart Tasche, Kay-Peter Hoyer, Evgeny Zhuravlev, Guido Grundmeier, Mirko Schaper and Olaf Keßler

485

Mechanical Behavior and Microstructure of Porous Ti Using TiC as Reinforcement . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Shiyuan Liu, Jian Wang, Tengfei Lu, Guibao Qiu and Hao Cui

495

Processing of Haynes® 282® Alloy by Laser Powder Bed Fusion Technology . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Robert Otto, Vegard Brøtan, Amin S. Azar and Olav Åsebø

503

Part IX

Advanced High-Strength Steels III

Tensile Deformation Behavior of 1 GPa-Grade TRIP-Aided Multi-microstructure Steels Studied by In Situ Neutron Diffraction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Noriyuki Tsuchida, Takaaki Tanaka and Yuki Toji Development of Advanced High-Strength Steels for Automobile Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Francys Barrado, Tihe Zhou, David Overby, Peter Badgley, Chris Martin-Root, Sarah Zhang and Rich Zhang Effect of Carbon Content on Strengthening Behavior with Grain Refinement on Lath Martensite Structure . . . . . . . . . . . . . . . . . . . . . Hiroyuki Kawata, Yoshiaki Honda and Kengo Takeda Assessment of the Strengthening Mechanisms Operating in Microalloyed Steels During Cyclic Deformation Using High-Resolution Electron Backscatter Diffraction . . . . . . . . . . . Paulina Lisiecka-Graca, Krzysztof Muszka and Janusz Majta

513

519

529

537

Effect of Niobium on Microstructure and Mechanical Properties of Nb–Ti Microalloyed Carbide-Free Bainitic Steels . . . . . . . . . . . . . Xi Chen, Fuming Wang, Changrong Li and Shuai Liu

549

Effect of Inclusions Modified by Y-Based Rare Earth on the Corrosion Behavior of EH36 Shipbuilding Steel . . . . . . . . . . . . . . . . . . . . . . . . Maolin Ye, Xiaojun Xi, Libin Zhu, Shufeng Yang and Jingshe Li

561

Microstructure and Mechanical Properties of Intercritical Annealed Multiphase Ultrahigh Strength Steel . . . . . . . . . . . . . . . . . . . . . . . . Huasai Liu, Xiangyu Li, Chunqian Xie and Yun Han

571

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xi

The Effect of Ni and Cu Addition on Mechanical Behavior of Thermomechanically Controlled Processed HSLA X100 Steels . . . . . . A. R. Hosseini Far, S. H. Mousavi Anijdan and M. Abbasi Part X

Advanced Magnetic Materials for Energy and Power Conversion Applications

Optimization of Magnetocaloric Properties of Ball-Milled La(Fe,Co,Si)13(H,C)y . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . V. Paul-Boncour, K. Nakouri and L. Bessais Production of High-Resistivity Electrical Steel Alloys by Substitution of Si with Al and Cr . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Brhayan Stiven Puentes Rodriguez, David Brice, James B. Mann, Srinivasan Chandrasekar and Kevin Trumble Nanocrystalline Multifunctional Pr–Co Compounds . . . . . . . . . . . . . W. Bouzidi, T. Bartoli, A. Michalowicz, J. Moscovici, N. Mliki and L. Bessais Part XI

579

593

599

607

Advanced Microelectronic Packaging, Emerging Interconnection Technology, and Pb-free Solder

A Study on Electrical Conductivity of Micro Friction Stir-Welded Dissimilar Sheets for Hybrid Electric Vehicles (HEVs) . . . . . . . . . . . . Omkar Mypati, Surjya Kanta Pal and Prakash Srirangam

619

Micro-structure and Properties of Cu–0.3 wt%Ag Alloy Ultra-Fine Wires . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Shu-sen Wang, Yuan-wang Zhang and Da-wei Yao

629

Length Scale of the Cellular Microstructure Tailoring Tensile Properties of Zn–20 wt%Sn–2 wt%Cu Solder Alloy . . . . . . . . . . . . . Cesar Bertolin dos Santos Mangualde, Rodrigo Valenzuela Reyes and José Eduardo Spinelli Effect of Ag on the Mechanical Properties of Bi–Ag Solder Alloys by the Single-Lap Shear Test Method . . . . . . . . . . . . . . . . . . . . . . . Nima Ghamarian, M. A. Azmah Hanim, M. Nahavandi, Ali Ourdjini, Zulkarnain Zainal and H. N. Lim Part XII

637

645

Advances in Computational Methods for Damage Mechanics and Failure Phenomena

Parametrically Homogenized Continuum Damage Mechanics (PHCDM) Models for Composites from Micromechanical Analysis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Xiaofan Zhang, Zhiye Li, Daniel J. O’Brien and Somnath Ghosh

657

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Contents

Effect of Multi-gating System on Solidification of Molten Metals in Spur Gear Casting: A Simulation Approach . . . . . . . . . . . . . . . . . Enesi Y. Salawu, Emuowhochere Oghenevwegba, Oluseyi O. Ajayi, A. O. Inegbenebor, E. T. Akinlabi and S. T. Akinlabi Part XIII

667

Advances in Surface Engineering

Corrosion Study of Boron Nitride Nanosheets Deposited on Copper Metal by Electrophoretic Deposition . . . . . . . . . . . . . . . . . . . . . . . . Mohsin Ali Raza, Amer Nadeem and Muhammad Tasaduq Ilyas Effects of Process Parameters on the Zirconia Coating Prepared by Sol-Gel and Electrodeposition Process . . . . . . . . . . . . . . . . . . . . . Jian Dong, Yanhui Sun, Bingsheng Dou, Feiyu He, Hongtao Huang and Jianping Zhen The Study of Slurry Erosion Wear Behaviour of Coal Bottom Ash Slurry Handling Pipeline . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Satish R. More, Sudeep P. Ingole, Dhananjay V. Bhatt and Jyoti V. Menghani Wear Characterization of Cemented Carbide Multipoint Cutting Tool Machining AISI 4140 at High Cutting Speed: Criteria for Materials Selection . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Federico Simone Gobber, Elisa Fracchia and Mario Rosso Dry Sheet Metal Forming Through Selective Oxidized Tool Surfaces . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Bernd-Arno Behrens, Deniz Yilkiran, Simon Schöler, Sven Hübner, Kai Möhwald and Fahrettin Özkaya

681

687

697

711

719

Effect of Process Parameters on Surface Properties of Laser-Hardened Cast Iron . . . . . . . . . . . . . . . . . . . . . . . . . . . . . S. V. Wagh, Sudeep Ingole, D. V. Bhatt, J. V. Menghani and M. J. Rathod

733

On Improvement in Surface Integrity of µ-EDMed Ti–6Al–4V Alloy by µ-ECM Process . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Ramver, Akshay Dvivedi and Pradeep Kumar

745

Corrosion and Wear Resistance of PTFE-Al2O3 Coatings Deposited on Aluminum Alloy by a Microblasting Process . . . . . . . . . . . . . . . . A. M. Oladoye, J. G. Carton, A. Baroutaji, M. Obeidi, J. Stokes, B. Twomey and A. G. Olabi

755

Contents

Part XIV

xiii

Algorithm Development in Materials Science and Engineering

Numerical Simulation of Ti6–Al4–V Alloy Diffusion Bonding Process Based on Molecular Dynamics . . . . . . . . . . . . . . . . . . . . . . . . . . . . Xiaogang Liu, Yongji Zuo and Haiding Guo Part XV

Alloys and Compounds for Thermoelectric and Solar Cell Applications VII

Custom Pyrolytic Graphite–Steel Thermocouple for High-Temperature Measurements . . . . . . . . . . . . . . . . . . . . . . . Abdul-Sommed Hadi and Bryce E. Hill Part XVI

765

781

Biological Materials Science

3D Contact and Strain in Alveolar Bone Under Tooth/Implant Loading . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Yuxiao Zhou, Chujie Gong, Mehran Hossaini-Zadeh and Jing Du Shear-Punch Testing of Human Cranial Bone and Surrogate Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . A. D. Brown, C. A. Gunnarsson, K. A. Rafaels, S. Alexander, T. A. Plaisted and T. Weerasooriya Investigation of Biodegradable Zn–Li–Cu Alloys for Orthopaedic and Cardiovascular Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . Jacob Young and Ramana G. Reddy Low-Temperature Air Plasma Modification of Electrospun Soft Materials and Bio-interfaces . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Bernabe S. Tucker, Ranu Surolia, Paul A. Baker, Yogesh Vohra, Veena Antony and Vinoy Thomas Accumulation of Biofilm on Ti–6Al–4V Alloy Fabricated Using Additive Layer Manufacturing . . . . . . . . . . . . . . . . . . . . . . . . . . . . Mari Koike, Tetsuro Horie, Richard J. Mitchell and Toru Okabe Copper Recovery from Printed Circuit Boards from Smartphones Through Bioleaching . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Lidiane Maria de Andrade, Carlos Gonzalo Alvarez Rosario, Mariana Alves de Carvalho, Denise Crocce Romano Espinosa and Jorge Alberto Soares Tenório Dependence of the Ferrovanadium Power as Additive on Mechanical Property in Porous Ti . . . . . . . . . . . . . . . . . . . . . . . Guibao Qiu, Jian Wang, Shiyuan Liu, Chenguang Bai and Yilong Liao

793

799

809

819

827

837

845

xiv

Contents

Effect of Compaction Pressure on Porosity and Mechanical Properties of Porous Titanium as Bone Substitute Materials . . . . . . . . . . . . . . . Qingjuan Li, Guibao Qiu, Shiyuan Liu and Tengfei Lu The Effect of Milling Time on Structural, Friction and Wear Behavior of Hot Isostatically Pressed Ti–Ni Alloys for Orthopedic Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Mamoun Fellah, Naouel Hezil, Mohammed Abdul Samad, Mohamed Zine Touhami, Alex Montagne, Alain Iost, Alberto Mejias and Stephania Kossman Part XVII

855

865

Bulk Metallic Glasses XVI

Perturbation Analysis of Amorphous Alloy Formation . . . . . . . . . . . . Rahul Basu

879

Shockwave Consolidation to Create Bulk Metallic Glass . . . . . . . . . . David Nemir, Jan Beck, Lawrence Murr, Yirong Lin and Luis Chavez

887

Part XVIII

Ceramic Materials for Nuclear Energy Research and Applications

Characterization of the Irradiation Effects in Nuclear Graphite . . . . . J. David Arregui-Mena, Philip D. Edmondson, Robert N. Worth, Cristian Contescu, Timothy D. Burchell and Yutai Katoh

901

Irradiation Effects on Reactor Concrete Structures . . . . . . . . . . . . . . J. David Arregui-Mena, Alain B. Giorla, G. E. Jellison, Elena Tajuelo-Rodriguez, Christa E. Torrence, Masaki Kawai, Yann Le Pape and Thomas M. Rosseel

907

Part XIX

Coatings and Surface Engineering for Environmental Protection

Electrochemical Mechanism and Preparation of Cr–Low-Carbon Steel Composite in a NaCl–KCl–NaF–Cr2O3 Molten Salt . . . . . . . . . . Shixian Zhang, Yungang Li, Cong Wang and Xiaoping Zhao

915

Diamond-Like Carbon Coating for Drill Collars: Test Experiences . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nausha Asrar and Jeffrey Ham

927

Inhibition Effect of Essential Oil Extracts on the Corrosion Inhibition of Mild Steel in Chloride–Sulphate Media . . . . . . . . . . . . . . . . . . . . Roland Tolulope Loto, Richard Leramo and Babatunde Oyebade

939

Contents

xv

Corrosion Properties of Steel Sheet with Zinc-Base Alloy Coatings . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Guangrui Jiang, Guanghui Liu, Ting Shang and Wanling Qiu Effect of Heat Treatment on the Localized Corrosion Resistance of S32101 Duplex Stainless Steel in Chloride/Sulphate Media . . . . . . . Roland Tolulope Loto, Cleophas Akintoye Loto, Akanji Olaitan and Olufunmilola Joseph Study of Mechanisms of Cobalt Electrodeposition by Means of Potentiodynamic Polarization Curves . . . . . . . . . . . . . . . . . . . . . . M. Ohba, T. Scarazzato, D. C. R. Espinosa, J. A. S. Tenório and Z. Panossian Part XX

959

967

Computational Approaches for Big Data, Artificial Intelligence and Uncertainty Quantification in Computational Materials Science

Artificial Intelligent and Simulation Nanostructure of Ceramic . . . . . . Habibollah Aminirastabi, Fatemeh Karimidehcheshmeh and Gouli Ji Part XXI

949

979

Computational Thermodynamics and Kinetics

Kinetics Calculation and Analysis of AlN Precipitation in ML40Cr Steel Austenite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Ziyi Liu, Yanping Bao and Min Wang

997

Study of Dendrite Growth Under Forced Convection in Superalloy Solidification by Multiphase-Field Coupled Lattice Boltzmann Method . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1007 Cong Yang, Qingyan Xu and Baicheng Liu Modeling of Volume Diffusion-Controlled Phase Transformations in Multiphase Multicomponent Alloy Systems by Minimization of Gibbs Energy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1019 Anders Salwén A New Method for Calculation of Vapor–Liquid Equilibrium (VLE) of Au–Cu Alloy System . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1027 Lingxin Kong, Jingbao Gao, Junjie Xu, Baoqiang Xu, Bin Yang and Yifu Li Ab Initio Study on the Oxidation Mechanism of Millerite . . . . . . . . . 1037 Xiaolu Xiong, Xionggang Lu, Guangshi Li, Hongwei Cheng, Qian Xu and Shenggang Li Kinetic Model of Silica Dissolution in CaO–SiO2–MgO–Al2O3 Slag System . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1045 Haifei An, Jie Li, Aimin Yang, Weixing Liu and Can Tian

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Part XXII

Contents

Deformation and Damage Behavior of High Temperature Alloys

Origin of the Significant Impact of Ta on the Creep Resistance of FeCrNi Alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1057 D. Magne, X. Sauvage and M. Couvrat Stress Analysis and Structure Optimization of W-Shaped Radiant Tube in Continuous Annealing Furnace . . . . . . . . . . . . . . . . . . . . . . 1067 Yang Long Li, Shun Ming Liu, Da Wei Hou, Wei Guo, Hui Wang and Meng Yu Part XXIII

Effective Business Improvement Methodologies for the Minerals, Metals, and Materials Industries

Case Studies of Continuous Improvement Projects in the Metals Industry . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1079 Cynthia Belt Part XXIV

Environmentally Assisted Cracking: Theory and Practice

Fracture Mechanics-Based Study of Stress Corrosion Cracking of SS304 Dry Storage Canister for Spent Nuclear Fuel . . . . . . . . . . . 1089 Leonardi Tjayadi, Nilesh Kumar and Korukonda L. Murty Similar and Dissimilar Metal Weld Failures in Hydrocracking Service at a Refinery . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1099 Sudhakar Mahajanam, Cesar Espinoza and Yenny Cubides Influence of Tempering Treatment on Precipitation Behavior, Microstructure, Dislocation Density and Hydrogen-Induced Ductility Loss in High-Vanadium Hot-Rolled X80 Pipeline Steel . . . . . . . . . . . 1111 Longfei Li, Bo Song, Zeyun Cai, Zhen Liu and Xiaokang Cui Part XXV

Fatigue in Materials: Multi-Scale and Multi-Environment Characterizations and Computational Modeling

Initiation and Early Growth of Fatigue Cracks . . . . . . . . . . . . . . . . . 1125 Jaroslav Polák Part XXVI Freeze Linings: Myth and Reality Influence of CaO/SiO2/Al2O3 Ratio on the Melting Behaviour of SynCon Slags . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1139 Dominik Hofer, Stefan Luidold, Tobias Beckmann and Frank Schulenburg

Contents

xvii

Freeze Lining Refractories in Non-ferrous TSL Smelting Systems . . . . 1149 Stanko Nikolic, Ben Hogg and Paul Voigt Freeze-Lining Formation in Submerged Arc Furnaces Producing Ferrochrome Alloy in South Africa . . . . . . . . . . . . . . . . . . . . . . . . . 1161 Joalet Dalene Steenkamp, Quinn Gareth Reynolds, Markus Wouter Erwee and Stefan Swanepoel Designing Furnace Lining/Cooling Systems to Operate with a Competent Freeze Lining . . . . . . . . . . . . . . . . . . . . . . . . . . . 1181 Hugo Joubert and Isobel Mc Dougall Part XXVII General Poster Session Corrosion Resistance of Hot Dipping Al–Zn–Si and Zn–Al–Mg–Si Alloy Coating . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1199 Hui Li, Jinglong Liang, Dongbin Wang and Yungang Li Performance of Low-Cost 3D Printed Pylon in Lower Limb Prosthetic Device . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1207 Fariborz Tavangarian, Camila Proano and Caleb Zolko Sequential Leaching Characteristics of Chromium in AOD Slag-Based Cementitious Materials . . . . . . . . . . . . . . . . . . . . . . . . . 1217 Ya-Jun Wang, Jun-Guo Li, Ya-Nan Zeng and Xiao-Yu Li Study on the Reaction Behavior of Hydrochloric-Acid-Containing Titanium Blast Furnace Slag . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1227 Jinglong Liang, Hui Li, Jing Wang, Dongbin Wang, Ramana G. Reddy and Yu Yang Thermodynamic and Kinetic Analysis of Inhomogeneous Distribution of Solute on Precipitations in as Cast Nb–V–Ti Microalloyed Steel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1237 Ya-Nan Zeng, Jun-Guo Li and Ya-Jun Wang Part XXVIII

Heterogeneous and Gradient Materials (HGM III): Tailoring Mechanical Incompatibility for Superior Properties

Radiation and Corrosion Resistances of 316LN Austenitic Stainless Steel by Rotationally Accelerated Shot Peening . . . . . . . . . . . . . . . . . 1251 Bin Yang, Xudong Chen, Yuntian Zhu and Yusheng Li

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Part XXIX

High Entropy Alloys VII

A Novel Dual-Phase Gradient Material of High-Entropy Alloy Prepared by Spark Plasma Sintering . . . . . . . . . . . . . . . . . . . . . . . . 1263 Wei Zhang, Mingyang Zhang, Fangzhou Liu, Yingbo Peng, Songhao Hu and Yong Liu Molecular Dynamics Simulations on the Mechanical Behavior of AlCoCrCu0.5FeNi High-Entropy Alloy Nanopillars . . . . . . . . . . . . 1271 Wei Li, Jing Tang, Qingyuan Wang and Haidong Fan Production of AlCoCrFeNiME-Based High-Entropy Alloys via Self-Propagating High-Temperature Synthesis . . . . . . . . . . . . . . . . . 1281 Murat Alkan, Esra Dokumaci, Berkay Türkoglu, Aslihan Kara, Büsra Aksu and Dilan Ugurluer Part XXX

ICME Education in Materials Science and Mechanical Engineering

ICME Applied in the Undergraduate Capstone Senior Design Sequence . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1291 Paul Sanders Part XXXI Interfaces in Structural Materials: An MPMD Symposium in Honor of Stephen M. Foiles Effect of a Vertical Twin Boundary on the Mechanical Property of Bicrystalline Copper Micropillars . . . . . . . . . . . . . . . . . . . . . . . . 1305 DeAn Wei, Haidong Fan, Jing Tang and Xu Zhang Part XXXII Materials for Molten Salt Energy Systems Effect of Ni on the Corrosion Behavior of Haynes 230 Alloy in MgCl2-KCl Salt . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1313 Yuxiang Peng and Ramana G. Reddy Part XXXIII

Mechanical Behavior of Nuclear Reactor Components

The Study of Mechanical Behaviour of Materials for the Nuclear Reactor Components in SUSEN Hot Cells . . . . . . . . . . . . . . . . . . . . 1325 Mariia Zimina, Petr Švrčula, Pavel Zháňal, Ondřej Libera, Stefan Zaunschirm and Ondřej Srba Investigation of Radiation Temperature and Straining Temperature Effects on the Screw Dislocation Mobility Evolution in Irradiated Ferritic Grains Using 3D Dislocation Dynamics . . . . . . . . . . . . . . . . . 1335 Yang Li, Christian Robertson, Xianfeng Ma and Biao Wang

Contents

Part XXXIV

xix

Mechanical Behavior Related to Interface Physics III

Mechanical Properties of Amorphous Silicon Nanoparticles . . . . . . . . 1347 D. Kilymis, C. Gerard and L. Pizzagalli Part XXXV

Modeling and Simulation of Composite Materials

Molecular Dynamics Simulation of the Structure and Transport Properties of xKF–yNaF–zAlF3 . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1357 Jie Li, Hui Guo, Hongliang Zhang, Ru Cai Li, Qiyu Wang, Jingkun Wang and Tianshuang Li Part XXXVI

Phase Stability, Phase Transformations, and Reactive Phase Formation in Electronic Materials XVIII

Microstructure Evolution and Physics Properties of Low Silver Copper Alloy Wires During In Situ Composite Preparation . . . . . . . . 1373 Yuan-wang Zhang, Shu-sen Wang and Da-wei Yao Part XXXVII

Phase Transformations and Microstructural Evolution

Heat Treatment Strategies to Improve the Quasi-Static and Dynamic Performance of Alpha + Beta Titanium Alloys . . . . . . . . . . . . . . . . . 1383 Alireza Fadavi Boostani, Shiraz Mujahid, Andrew L. Oppedal, Cory Krivanec, Wilburn R. Whittington, Paul G. Allison, Jishnu J. Bhattacharyya, Sean Agnew and Haitham El Kadiri Effect of Silicon Content on the Dilatometric Behavior of a Medium-Carbon Steel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1389 A. I. Gallegos Pérez, O. Vázquez Gómez, J. J. López Soria, H. J. Vergara Hernández and E. López Martínez Phase-Field Simulation of Intermetallic Phase Precipitation in a High-Al Alloyed Lightweight High-Strength Steel . . . . . . . . . . . . 1401 Carsten Drouven, Bowen Zou, Wenwen Song and Wolfgang Bleck Shape Memory Behavior of Ni49.5Ti50.5 Processing-Induced Strain Glass Alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1411 Robert W. Wheeler, Jesse Smith, Nathan A. Ley, Anit Giri and Marcus L. Young Precipitation Hardening of Supersaturated Al–Sc–Zr Produced via Melt-Spinning . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1421 Yang Yang and Paul Sanders

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Contents

Effect of Sm Content and Solidification Rate on Microstructure of SmFe Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1427 Kun Liu, Shuhuan Wang, Yunli Fen, Chunyan Song, Guolong Ni and Kaixuan Zhang Evolution of Dendritic Morphology Under HPMO Treatment . . . . . . 1437 Hui-cheng Li, Yu-xiang Liu, Zhen Liu and Qi-jie Zhai In Situ Observation of Melting and d $ c Phase Transformation in Duplex Stainless Steel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1447 Yang Liu and Yan-hui Sun Microstructural Evolution of a Transformation in Which There Is an Exclusion Zone Around Each Nucleus . . . . . . . . . . . . . . . . . . . 1459 Paulo R. Rios, Harison S. Ventura, André L. M. Alves, Weslley L. S. Assis and Elena Villa Thermodynamic Properties of Si–B Alloys Determined by Solid-State Heterogeneous Phase Equilibrium . . . . . . . . . . . . . . . . . . . . . . . . . . 1471 Muhammad A. Imam and Ramana G. Reddy Part XXXVIII

Powder Processing of Bulk Nanostructured Materials

Microstructure Evolution and Mechanical Properties of Medical Material Mg–3Zn Alloy Prepared by Semi-solid Powder Injection Moulding . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1483 Xia Luo, Chao Fang, Zhou Fan, Bensheng Huang and Jun Yang Inhomogeneity of Strain in Metal Particulates Produced by Modulation-Assisted Machining . . . . . . . . . . . . . . . . . . . . . . . . . 1499 Indrani Biswas, James B. Mann, Srinivasan Chandrasekar and Kevin Trumble Numerical Simulation and Validation of Gas and Molten Metal Flows in Close-Coupled Gas Atomization . . . . . . . . . . . . . . . . 1507 F. Hernandez, T. Riedemann, J. Tiarks, B. Kong, J. D. Regele, T. Ward and I. E. Anderson Density Separation of Mixed Carbide Colloids via Standing Wave Physics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1521 Trenin K. Bayless, Jerome P. Downey, Grant C. Wallace and Mark D’Aberle The Influence of Mechanical Activation on the Synthesis of Ca2MgSi2O7 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1533 Fariborz Tavangarian and Caleb Zolko

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Part XXXIX

Recent Advances in Functional Materials and 2D/3D Processing for Sensors and Electronic Applications

3D Printing of Polymer-Based Gasochromic, Thermochromic and Piezochromic Sensors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1545 Patrick Dzisah, Airefetalo Sadoh and Nuggehalli M. Ravindra 3D Printing of Pharmaceuticals and Transdermal Drug Delivery––An Overview . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1563 David Bird, Emel Eker and Nuggehalli M. Ravindra Formulation of Curable Resins Utilized in Stereolithography . . . . . . . 1575 David Bird, Elbert Caravaca, Joseph Laquidara, Keith Luhmann and Nuggehalli M. Ravindra MARS––Magnetic Augmented Rotation System . . . . . . . . . . . . . . . . 1589 Vishwas Danthi Shivaram, Roulei Liu, Navjot Panchhi, Laila Alqarni, Rayan Daroowalla, Shuang Du, Yan Liu, Tien See Chow and Nuggehalli M. Ravindra Part XL

Recent Developments in Biological, Structural and Functional Thin Films and Coatings

Friction Conditions on Deep-Drawing Tool Radii When Using Volatile Media as Lubrication Substitute . . . . . . . . . . . . . . . . . . . . . . . . . . . 1603 Gerd Reichardt and Mathias Liewald Investigation of Friction and Adhesion Behavior of Textured Workpieces and Coated Tools Under Dry Tribological Contact . . . . . 1615 Rafael Hild, Robby Mannens, Daniel Trauth, Patrick Mattfeld, Thomas Bergs, Dennis C. Hoffmann, Nathan C. Kruppe, Tobias Brögelmann and Kirsten Bobzin Effects of Emissivity on Combustion Behavior of Energetic Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1629 Elbert Caravaca, David Bird, Henry Grau, Viral Panchal and Nuggehalli M. Ravindra Self-healing in Materials: An Overview . . . . . . . . . . . . . . . . . . . . . . 1643 Samiha Hossain and Nuggehalli M. Ravindra Part XLI

Solidification Processing of Light Metals and Alloys: An MPMD Symposium in Honor of David StJohn

Revealing the Heterogeneous Nucleation and Growth Behaviour of Grains in Inoculated Aluminium Alloys During Solidification . . . . . 1665 Yijiang Xu, Daniele Casari, Ragnvald H. Mathiesen and Yanjun Li

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Influence of Microstructure Evolution During Twin-Roll Casting on the Properties of Magnesium Sheets . . . . . . . . . . . . . . . . . . . . . . 1677 K. U. Kainer, G. Kurz, S. Pakulat and D. Letzig A History of the Global Light Metals Alliance . . . . . . . . . . . . . . . . . 1687 Jennifer Jackman, Kumar Sadayappan and Mark Easton Analysis of the High-Purity Aluminum Purification Process Using Zone-Refining Technique . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1697 Heli Wan, Baoqiang Xu, Jinyang Zhao, Bin Yang and Yongnian Dai Author Index . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1707 Subject Index. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1715

Part I

2019 International Metallurgical Processes Workshop for Young Scholars (IMPROWYS 2019)

Ab Initio Molecular Dynamics Study on the Dissolution of Interfacial Iron Oxides in Hot Compressive Bonding Combined with Experiments Honglin Zhang, Mingyue Sun, Bin Xu and Dianzhong Li

Abstract Based on experimental observation, ab initio molecular dynamics was used to investigate the dissolution of interfacial iron oxides in hot compressive bonding (HCB). The surface analysis indicated that there was reoxidation at the unclosed iron surface during the sample heating in HCB. The bonding of pre-oxidized iron was designed to verify the dissolution of iron oxides into matrix. Two models were proposed to understand the dissolution behavior with dynamic simulations. Model I was applied to the case with bonding interface between oxides and matrix,in which periodical interface structure of Fe3 O4 /BCC-Fe was constructed. The dissolution of Fe3 O4 contained the initial structural dissociation and the diffusion of free oxygen and iron atoms into matrix. It was found that the diffusivity of iron was higher than oxygen. Model II with embedded structure of oxide cluster was proposed to understand the initial dissolution of iron oxide particles in the matrix. The mean square displacement (MSD) results suggested that the local strain may promote the process by increasing the mobility of oxygen. And the Bader charge analysis implied that the electron contribution of iron matrix and its transfer to the dissociated atoms plays a key role in the initial dissolution of interfacial iron oxides. Keywords Hot compressive bonding · Interfacial oxides · Dissolution behavior · Ab initio molecular dynamics

H. Zhang · M. Sun · B. Xu · D. Li Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China H. Zhang School of Materials Science and Engineering, University of Science and Technology of China, Hefei 230026, China M. Sun (B) · B. Xu Key Laboratory of Nuclear Materials and Safety Assessment, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China e-mail: [email protected] © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_1

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Introduction As a conventional joining technology, diffusion bonding (DB) has been widely used in industrial manufacturing [1, 2]. Recently, based on the principles of DB, hot compressive bonding (HCB) has been used in joining metals such as steels and super alloys [3, 4]. Although there is an application of plastic deformation at elevated temperature (usually higher than DB), some adverse factors still imped the metallurgical bonding process, especially, for the oxide scales on the metal surface [5]. Normally, in addition to the native oxide layer, the reoxidation of bonding surface may happen during the heating stage [6]. Subsequently, the oxides existing at the bonding interface are regarded as a barrier of atomic diffusion and the migration of dislocations [5]. Therefore, the evolution of interfacial oxides and its impact on the bonding strength have been widely investigated. Zhu et al. [7] investigated the effect of oxidation level on the shear property of hot-rolled stainless steel clad plate, found that the low-oxidized cladding interface was important for obtaining the highproperty clad plates. Koyama et al. [8] observed that interfacial fine oxide of tin disappeared and granular oxide coalesced with the rise of bonding temperature and pressure, which obviously increased the joint strength. Recently, Sridharan et al. [9] directly detected the elements distribution at Fe–Al bonding interface through atom probe tomography, and verified oxide dissolution through the observation of non-equilibrium colossal super saturation of oxygen. Comparably, there were few studies focusing on the removal mechanism of interfacial oxides in diffusion bonding. Takahashi et al. [10] presented two diffusion models to explain the dissolution process of surface oxide on copper and titanium. Besides, reduction mechanism was proposed with respect to aluminium alloys [11], wherein the stable films of Al2 O3 could react with magnesium and transform to fine particles of Al2 MgO4 . However, to the best of our knowledge, there is a lack of investigation of the interaction and mass transfer between interfacial oxides and matrix at the atomic scale. The experimental characterization and dynamic simulation are combined to investigate the dissolution behavior of iron oxide at the HCB interface. Firstly, the surface film of pure iron was identified by X-ray photoelectron spectroscopy (XPS) and X-ray diffraction (XRD). Then the evolution of interfacial oxides in the HCB specimen was observed by optical microscope (OM). Furthermore, ab initio molecular dynamics (AIMD) based on density functionals theory (DFT) is adopted to simulate the initial evolution steps of the oxide and its reaction mechanism with matrix. Two models are given to facilitate the understanding of how interfacial oxides are removed.

Methodology The material in this work was pure iron with the chemical composition of 0.003%C–0.04%Mn–0.002%P–0.002%S. After cutting, mechanical fraying, and polishing; thin iron foils with thickness of 0.5 mm were oxidized for 48 h in air.

Ab Initio Molecular Dynamics Study on the Dissolution …

5

Fig. 1 Schematic diagrams: a specimens of HCB-1, b specimens of HCB-2, c HCB process, d process routes of HCB-1 and HCB-2

Then the foils were subjected to isothermal heat treatment at 1200 °C for 1 h in vacuum with 0.1–0.01 torr. And the surface analyses were implemented on Escab-250 XPS spectrometer in the constant analyser mode at pass energy of 100 eV with a step size of 1 eV, while the parameters for high resolution spectra were selected as 50 eV and 0.1 eV, respectively. The obtained high-resolution XPS spectra were fitted using the software XPSPEAK 4.1. In HCB experiments, first, the normal bonding specimens (Fig. 1a) were mechanically frayed to remove the oxide scales. Meanwhile, the specimens with rectangular notch were machined as illustrated in Fig. 1b. And they were preoxidized in furnace at 200 °C for 2 h. Then the bonding surface was frayed except for the notch. The type of surface oxide was identified by XRD (Cu Kα radiation, step size: 0.04°, scan rang: 28°–68°). Subsequently, both of them were hold in Gleeble-3500 thermal simulator as depicted in Fig. 1c. The given vacuum degree was in the range of 0.1–0.01 torr. The HCB process routes in detail were given in Fig. 1d, HCB-1 was implemented to join the normal specimens in order to observe the typical interfacial microstructure. In addition, HCB-2 of preoxidized specimens was designed to availably observe the evolution of interfacial iron oxides. The final bonded samples were cut along the direction vertical to the interface. After being mechanically polished and etched with 10 vol. % nital, the interfacial microstructure was observed on Axio Vert. A1 microscope, and the thickness of oxide film was statistically measured. AIMD simulations were carried out by using Vienna Ab Initio Simulation Package (VASP) [12]. The generalized gradient approximation (GGA) within the Perdew–Burke–Ernzerhof (PBE) parameterization scheme for the exchange–correlation function was adopted. To save the computational cost, a Monkhorst–Pack grid of 1 × 1×1 k points sampling and plane-wave cutoff energy of 400 eV were selected. And the time step was set as 1.0 fs within the canonical (NVT) ensemble at a constant temperature of 1500 K controlled by Nose–Hoover thermostats [13, 14]. Based on the experimental results, two interfacial oxides models were proposed and used to

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simulate the evolution process of interfacial iron oxides to understand the dissolution mechanism. The details about model construction were given in the following section.

Results and Discussion The Analysis of Surface Film In HCB process, the vacuum degree of the ambient environment plays a key role in the removal of surface oxide film. Figure 2 shows the depth profile of the composition of surface film under different conditions. The content of oxygen is higher after the heat treatment at 1200 °C in vacuum. Besides, the content of oxygen rapidly decreases from 70 at. % to 20 at. % before 100 s, and then its descending rate slows down. Based on the etching rate of 0.1 nm/s, the thickness of the film is about 10 nm, which is in agreement with the results reported by Sewell [15]. Furthermore, the curve fitting of the Fe2p peak in the film is presented in Fig. 3 to identify the type and content of surface oxides. It should be noted that since the reduction effect of Ar+ , the proportion of high-valence iron oxide may be underestimated. Here only the outmost surface and the layer of 2 nm in depth were measured. The resulting fit suggests there is a mix of iron, Fe3 O4, and Fe2 O3 in the outmost surface (Fig. 3a), while there are primarily Fe3 O4 and iron peaks when the surface is further etched to the depth of 2 nm (Fig. 3b). After heat treatment in vacuum, the content of Fe3 O4 increases from 52.1% to 75.2% as shown in Fig. 3c. And more iron phase is identified at the layer of 2 nm in depth (Fig. 3d). The above result implies that the iron surface can be reoxidized in the present vacuum degree. The oxygen partial pressure under the condition allows the formation of FeO at 1200 °C. And it can be transformed to Fe3 O4 when the temperature is below 570 °C based on the Fe–O phase diagram, which leads to the higher content of Fe3 O4 . Based on that, considering the same given vacuum degree in HCB process, the unclosed iron surface can be reoxidized during the heating period.

HCB Interfacial Microstructure of Pure Iron In Fig. 4a, there are fine and uniform equiaxed grains the matrix of as-received iron. After HCB-1 process without holding, the bonding interface can be found at the center of the sample, and there are elongated inhomogeneous grains distributed near the interface, as shown in Fig. 4b. And a part of grains migrates across the interface, leaving the microvoids wrapped inside them. This phenomenon indicates that there is dynamic recrystallization (DRX) in the matrix on the both sides of the interface

Ab Initio Molecular Dynamics Study on the Dissolution …

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Fig. 2 Depth profile of the main composition elements for the iron surface film

Fig. 3 Fe2p high resolution fitted peaks for: a the utmost surface in native condition, b the layer of 2 nm in depth in native condition, c the utmost surface after heat treatment, d the layer of 2 nm in depth after heat treatment

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Fig. 4 Optical micrographs shows the sample microstructure of: a as-received state, b after HCB-1 process, c after HCB-2 without holding, d after HCB-2 with holding of 1 h, e after HCB-2 with holding of 3 h, f after HCB-2 with holding of 6 h

[16]. The voids can gradually vanish with the volume diffusion of atoms [5], and the final sound metallurgical bonding is achieved. Based on the analysis of the surface film and the microstructure of HCB-1, it is difficult to experimentally observe the evolution of iron oxides. Therefore, the following presents the results of HCB-2 with preoxidized specimens. Figure 4c shows the optical microstructure of the sandwich structure of matrix/oxide film/matrix after HCB-2 process. There is an obvious crack inside the oxide film due to the deformation stress, and the broken oxides particles can also be observed. Due to the obstacle of oxide film, only a few matrix grains extrude along the above crack. After holding at 1200 °C for 1 h (Fig. 4d), the interfaces gradually become straight, while the crack still exists. And newly formed fine grains appear at the upper interface and the grain extrusion becomes obvious at the lower interface with the time (Fig. 4e). After holding for 6 h, the average thickness of oxide film decreases from 11.9 to 9.6 μm, and it turns to be discontinuous as shown in Fig. 4f. Although, the oxide film does not disappear, the above result suggests the dissolution of the interfacial oxide film at 1200 °C. In addition, as shown in Fig. 5, the XRD results indicate the main oxide phase of the film is Fe3 O4 after the surface being preoxidized for 2 h. Furthermore, based on the above experimental characterization, AIMD simulations have been carried out as follows to investigate the dissolution of interfacial iron oxides at the atomic scale.

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Fig. 5 XRD pattern of iron surface with preoxidation for 2h

Interfacial Oxides Dissolution by AIMD Based on the experimental observations, there are two forms of interfacial oxides: (1) the oxide film bonded with matrix, (2) broken oxide particles wrapped in matrix. Taking the computational cost into consideration, two models of bonding interface (Model I) and oxide cluster embedded (Model II) are proposed to simulate the initial evolution process of Fe3 O4 in body-centered cubic (BCC) iron matrix. The bonding interface model is given first. 1. Model I The interface structure of Fe3 O4 (001)/BCC-Fe (001) with orientation relationship of Fe3 O4 [110]/BCC-Fe [100] is constructed based on the lattice parameters of Fe3 O4 (a  8.40 Å) and BCC-Fe (a  2.84 Å). The structure is more stable when the FeO2 terminal of Fe3 O4 occupies at the interface [17]. The upper and lower interface distances are set as 1.5 Å and 2 Å, respectively. Due to the existence of defects inside the oxide film, the lattice parameter of BCC-Fe is kept to deal with the mismatch. The evolution of the configuration of Fe3 O4 /BCC-Fe interface is shown in Fig. 6. Initially, the bonds of Fe–O break with the dissociation of structure, and then part of dissociated oxygen and iron atoms in Fe3 O4 are prone to segregate at the interface. With the evolution progress, more oxygen and iron atoms dissociated from the oxide move across the interfaces at both sides and permeate into the matrix (as the blue arrows indicate). And it can be found that the migration distance of iron is larger than oxygen, which is related with their diffusivity. The dissociated iron atoms can enter the matrix more easily by substitutional diffusion mechanism, especially if there is high occurrence possibility of vacancies by thermal activation. By contrast, due to the interaction with vacancies in BCC-Fe [18], oxygen atoms tend to be trapped, leading to a short diffusion distance. After 16 ps, the above diffusion becomes more obvious with iron atoms of matrix diffusing into the oxides (as the green arrows indicate). And the previous interface gradually disappears with the time.

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Fig. 6 A series of snapshots of bonding interface model at various evolution time, red and orange balls are oxygen and iron atoms, respectively, in Fe3 O4 , pink balls indicate the iron atoms in matrix, the diffusion direction of dissociated atoms (blue arrow), the iron atoms diffused from matrix (green arrow) (Color figure online)

The oxygen mean displacement along z-direction (vertical to the interface) z d (t) is given in Fig. 7. Before 9 ps, the diffusion distance of oxygen at the lower interface linearly increases and then fluctuates around 2 Å (Fig. 7a), indicating the oxygen atoms migrates near the interface. After 15 ps, the value exceeds the interface distance, and the oxygen diffuses into the matrix. As shown in Fig. 7b, the diffusion behavior of oxygen near the upper interface is similar during the evolution process. Additionally, Fig. 7c–d shows the maximum value of 6 Å along z-direction for oxygen at both interfaces, implying the oxygen exit in the matrix. Based on the analysis of the atomic configuration and displacement, the dissociation of Fe3 O4 and further diffusion process are simulated, which facilitates to understand the dissolution process of oxide film. 2. Model II Previous studies have showed that the oxide film of copper could be vaporized by dispersing; globularizing and dissolution in sequence [10], and the interfacial irregular oxides could be gradually turned into spherical shape before vanishing during the post-holding treatment [19]. The morphology evolution of interfacial oxides indicates the dissolution process initiates at the surface metastable region. Therefore, Fe3 O4 cluster is assumed as a metastable phase that is embedded in the BCC-Fe matrix to simulate the initial process of dissolution. Based on the previous method [20], random two layers of Fe3 O4 along z-axis are selected and noted as Fe4 O5 . Meanwhile, iron atoms are removed from the BCC-Fe supercell to create vacancies, and they are filled by the cluster of Fe4 O5 . The model is depicted as the initial configuration in Fig. 8. If there is no strain applied, the cluster rapidly dissociates into free oxygen and iron atoms within 2 ps, following by the migration into the matrix with the time. After 10 ps, the embedded cluster completely dissolves into the matrix with the interdiffusion of iron atoms in matrix. In order to investigate the effect of

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Fig. 7 Mean displacement along z-direction: a oxygen atoms at the lower interface, b oxygen atom with the maximum value at the lower interface, c oxygen atoms at the upper interface, d oxygen atom with the maximum value at the upper interface, the dashed line indicates the interface distance

strain, first, 5% compressive strain of BCC-Fe lattice is introduced near the cluster as shown in Fig. 8b. Although the dissociation process is more obvious before 5 ps, the evolution process is basically unchanged during the whole time. Furthermore, after applying 2% x-y plane tensile strain on the lattice of iron matrix (Fig. 8c), the interaction is more obvious near the cluster as before, while the configuration of iron atoms in matrix is basically unchanged. It may be related to the rising potential of atoms caused by the lattice stretching. In order to further evaluate the effect of strain on the mobility of oxygen, Fig. 9a shows the MSD for the three cases. It can be seen that all values gradually increase with the time, indicating the oxygen atoms gradually migrate away from their initial position. After applying 5% compressive strain along z-axis, the MSD value increases obviously compared with the no strain applied. However, the triaxial strain does not further increase the MSD value. As for the mean displacement along z-direction (Fig. 9b), the effect of triaxial strain is more obvious before 7 ps. It may attribute to that more space in the lattice generated by the x-y plane tensile allows oxygen atoms to move faster along z-direction in the primary period. With the evolution progress, the diffusion distance of oxygen for both strain-applied cases can reach to the same

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Fig. 8 A series of snapshots of cluster embedded model at various evolution time considering: a without strain, b with strain along z-direction, c with triaxial strain

level. The above results suggest that strain may facilitate the dissolution process of oxide particles with respect to the diffusion of oxygen. Bader charge analysis is used to understand the interaction mechanism between iron oxide cluster and matrix. Bader charge is related to the valence electrons of selected potentials. Typically, the differential between it and valence charge indicates the valence state [21]. Figure 10 gives the evolution of the Bader charge for all atoms in the model. Initially, in Fig. 10a, the Bader charge of oxygen atoms Q O is around 6.8, and the Bader charge of iron atoms in oxide cluster Q cFe depends on their sites with the minimum value of 7.4. Besides, the bader charge of iron atoms in matrix Q mFe

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Fig. 9 The mobility of oxygen evaluated by: a MSD value, b mean displacement value along z-direction

is around 8.0, while the charge becomes obviously disperse for the iron atoms near the cluster. At 0.6 ps, both of Q O and Q cFe somewhat increase (Fig. 10b). Comparably, the value of Q mFe for these iron atoms near the cluster decreases to 7.4 and less. It implies there is electrons transfer between the atoms in oxide cluster and matrix. With the time increases, as shown in Fig. 10c, the value of Q O is within 6.9–7.0, while the average value of Q cFe further increases and reaches to 8.0. Furthermore, when the time is up to 4 ps (Fig. 10d), the value of Q O is around 7.0, indicating the valence state of the oxygen atom is basically unchanged. And Q cFe distributes within the range of Q mFe . It suggests the dissociated iron atoms to be identical with the iron atoms of matrix. Based on the above results, there is electrons transfer between the iron oxide cluster and matrix. On the one hand, such transfer may promote the formation of a covalent bond between oxygen and transition metal (like iron in the matrix) based on previous studies [22]. Besides, the dissociated iron can receive electrons and changes to free metal atoms from previous the ionic state. Therefore, the electron contribution of iron matrix and its exchange with oxide cluster accelerate the initial dissolution, especially for the dissociation process.

Conclusions The experimental characterization and AIMD simulations were performed to understand the dissolution mechanism of interfacial oxides in the HCB of pure iron. After preoxidation, the oxide film with thickness of 10 nm is mainly consisted of Fe3 O4 . And the interfacial iron oxides could hardly be observed due to the DRX. The middle iron oxide layer became thinner and discontinuous because of its dissolution into matrix as well as the growth and extrusion of grains. The bonding interface model showed the dissociated oxygen and iron atoms could further diffuse into BCC-Fe matrix with different distance due to various diffusion mechanisms. And the

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Fig. 10 Evolution of bader charge at the time of: a 0 ps, b 0.6 ps, c 2 ps, d 4 ps

maximum diffusion distance of oxygen atoms exceeded the lattice parameter of BCCFe. Additionally, the oxide cluster embedded model indicated the local strain may promote the initial dissolution of oxide particle by increasing the mobility of oxygen. The dissociated oxygen atoms achieved electrons and formed covalent bonds with iron atoms in matrix. And electrons also transferred to the dissociated iron to promote the formation of free atoms. This process accelerated the initial dissolution of the oxide particle in matrix. Acknowledgements We thank X. Q. Chen for theoretical assistance and valuable discussions. All calculations have been performed on the high-performance computational cluster in the Shenyang National University Science and Technology Park. The authors are also grateful to the financial support from National Key Research and Development program (Grant No. 2016YFB0300401), National Natural Science Foundation of China (Grant Nos. U1508215, 51774265) and Key Program of the Chinese Academy of Sciences (Grant No. ZDRW-CN-2017-1).

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References 1. Han WB, Zhang KF, Wang GF (2007) Superplastic forming and diffusion bonding for honeycomb structure of Ti-6Al-4 V alloy. J Mater Process Technol 183(2–3):450–454 2. Noh S, Kasada R, Kimura A (2011) Solid-state diffusion bonding of high-Cr ODS ferritic steel. Acta Mater 59(8):3196–3204 3. Gao XJ, Jiang ZY, Wei DB, Jiao SH, Chen DF, Xu JZ, Zhang XM, Gong DY (2014) Effects of temperature and strain rate on microstructure and mechanical properties of high chromium cast iron/low carbon steel bimetal prepared by hot diffusion-compression bonding. Mater Des 63:650–657 4. Yang XW, Li WY, Feng Y, Yu SQ, Xiao B (2016) Physical simulation of interfacial microstructure evolution for hot compression bonding behavior in linear friction welded joints of GH4169 superalloy. Mater Des 104:436–452 5. Kazakov NF (1985) Diffusion bonding of materials. Pergamon, Oxford 6. Tsukamoto M (2016) Improvement of diffusion bondability on chromium-copper alloy by behavior of oxide layer at high temperature. J Jpn Inst Met 80(3):206–212 7. Zhu ZC, He Y, Zhang XJ, Liu HY, Li X (2016) Effect of interface oxides on shear properties of hot-rolled stainless steel clad plate. Mater Sci Eng A-Struct Mater Prop Microstruct Process 669:344–349 8. Koyama S, Takahashi M, Ikeuchi K (2004) Behavior of superficial oxide film at solid-state diffusion-bonded interface of tin. Mater Trans 45(2):300–302 9. Sridharan N, Isheim D, Seidman DN, Babu SS (2017) Colossal super saturation of oxygen at the iron-aluminum interfaces fabricated using solid state welding. Scr Mater 130:196–199 10. Takahashi Y, Nakamura T, Nishiguchi K (1992) Dissolution process of surface oxide film during diffusion bonding of metals. J Mater Sci 27(2):485–498 11. Ikeuchi K, Matsuda F, Kotani K (1996) Behaviour of oxide at diffusion-bonded interfaces in Al-Mg-Si series 6063 alloy: Study of diffusion-bonding mechanism of aluminium alloys by transmission electron microscopy (1st Report). Weld Int 10(9):697–704 12. Kresse G, Furthmuller J (1996) Efficient iterative schemes for ab initio total-energy calculations using a plane-wave basis set. Phys. Rev. B 54(16):11169–11186 13. Hoover WG (1985) Canonical dynamics-equilibrium phase-space distributions. Phys Rev A 31(3):1695–1697 14. Nose S (1984) A molecular dynamics method for simulations in the canonical ensemble. Mol Phys 52(2):255–268 15. Sewell PB, Stockbridge CD, Cohen M (1961) An electrometric and electron diffraction study of air-formed oxide films on iron. J Electrochem Soc 108(10):933–941 16. Zhang JY, Sun MY, Xu B, Li DZ (2018) Interfacial microstructural evolution and metallurgical bonding mechanisms for IN718 superalloy joint produced by hot compressive bonding. Metall Mater Trans B-Proc Metall Mater Proc Sci 49(5):2152–2162 17. Forti MD, Alonso PR, Gargano PH, Balbuena PB, Rubiolo GH (2016) A DFT study of atomic structure and adhesion at the Fe(BCC)/Fe3O4 interfaces. Surf Sci 647:55–65 18. Barouh C, Schuler T, Fu CC, Jourdan T (2015) Predicting vacancy-mediated diffusion of interstitial solutes in alpha-Fe. Phys Rev B 92(10):104102 19. Xie B, Sun M, Xu B, Wang C, Li D, Li Y (2018) Dissolution and evolution of interfacial oxides improving the mechanical properties of solid state bonding joints. Mater Des 157:437–446 20. Danielson T, Tea E, Hin C (2016) Investigation of helium at a Y2Ti2O7 nanocluster embedded in a BCC Fe matrix. Phys Chem Chem Phys 18(43):30128–30134 21. Cen WL, Liu Y, Wu ZB, Wang HQ, Weng XL (2012) A theoretic insight into the catalytic activity promotion of CeO2 surfaces by Mn doping. Phys Chem Chem Phys 14(16):5769–5777 22. Matsunaga K, Sasaki T, Shibata N, Mizoguchi T, Yamamoto T, Ikuhara Y (2006) Bonding nature of metal/oxide incoherent interfaces by first-principles calculations. Phys Rev B 74(12):125423

Effect of MgO Content on the Properties of Magnesia Fluxed Pellets Yuzhu Zhang, Weixing Liu, Aimin Yang and Jie Li

Abstract With the increasing shortage of iron concentrate powder resources, the optimization of blast furnace charge structure by using the fluxed pellets has become an important method to reduce the cost of iron making in iron and steel enterprises. To solve some problems, such as bonding, ringing, and so on; the effect of MgO on the properties of pellets has been studied. The results show that with the increase of MgO, the compressive strength of green pellets increased, while the dropping strength of green pellets decreased. The decrepitation temperature increased at first and then decreased, and the bond rate decreased. The performance of the pellets improved. The hematite content decreased and the magnesium ferrite content increased in main crystal phase of the pellets, which was mainly linked by crystal bond. Keywords Fluxed pellets · MgO · Basicity · Microstructure

Introduction With the continuous exploitation of earth resources and a shortage of resources, the optimization of blast furnace burden structure has become an important means for steel enterprises to reduce the cost of iron smelting. The poor performance of ordinary acid pellets makes it a serious constraint on the improvement of the ratio of entering furnace. Fluxing pellets can improve the properties of pellets, but it is easy to cause problems such as bonding and knots. By adding MgO, high-temperature bonding can be avoided, and also the performance of the pellet can be improved, so the focus of domestic and foreign scholars began to shift to magnesium fluxed pellets. Fan et al. [1], Pan et al. [2], Li et al. [3], Gao et al. [4, 5], Qing et al. [6, 7], Sui et al. [8] have a lot of study results different magnesium fluxes of effects to magnesium fluxed pellets of properties and their mechanism. Cooke and Brandt [9], Srinivas et al. [10, 11], respectively, from the effects of magnesium fluxed pellets Y. Zhang · W. Liu (B) · A. Yang · J. Li North China University of Science and Technology, Tangshan City, China e-mail: [email protected] © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_2

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on the properties, microstructure and mechanism of action of pellets were studied. This article mainly starts with the raw material structure, the effect of MgO pellet explosion temperature was studied. It provides theoretical and technical support for industrial production and application of magnesium-fluidized pellets.

Experimental Materials and Methods Experimental Materials Raw Material Chemical Composition The chemical composition of mineral powder and flux selected in the experiment and physical properties of bentonite are shown in Tables 1 and 2.

Table 1 Chemical composition of raw materials/% Raw material

TFe

FeO

SiO2

CaO

MgO

Al2 O3

Iron ore powder 1

66.56

27.98

5.46

Iron ore powder 2

64.33

26.62

1.32

Iron ore powder 3

66.45

24.56

6.31

Iron ore powder 4

47.33

0.061

2.04

Limestone

–

–

3.22

49.92

Light burning magnesium powder

–

–

5.92

Bentonite

–

–

56.23

RI

0.25

0.56

1.38

−3.06

0.91

3.33

0.92

−1.60

0.15

0.26

0.84

−2.48

0.079

0.54

9.94

14.34

0.54

1.16

44.04

1.70

83.64

0.99

6.96

4.59

2.56

12.12

11.92

Raw material

K2 O

Na2 O

Zn

Mn

Cu

Cr

Ni

Iron ore powder 1

0.055

0.047

0.0051

0.052

0.0018

0.0052

0.0016

Iron ore powder 2

0.046

0.014

0.024

0.18

0.064

0.020

0.028

Iron ore powder 3

0.074

0.012

0.0010

0.058

0.0011

0.024

0.0012

Iron ore powder 4

0.022

0.010

0.034

0.77

0.018

0.18

0.51

Table 2 Physical properties of bentonite Raw material

Colloid valence/%

Expansion capacity/(mL g−1 )

Blue uptake/g

Montmorillonite content/%

Bentonite

218.00

9.00

28.31

64.04

Effect of MgO Content on the Properties of Magnesia … Table 3 Phase composition of iron ore powder/%

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Raw material

MFe

CFe

SFe

OFe

SiFe

Iron ore powder 1

65.46

0.085

0.030

0.44

0.25

Iron ore powder 2

63.10

0.14

0.048

0.58

0.11

Iron ore powder 3

63.62

0.17

0.34

1.93

0.40

Iron ore powder 4

0.17

0.11

0.12

46.36

0.54

Composition of Iron Ore Powder Phase Table 3 shows that iron ore powder 1, 2 and 3 are mainly magnetite; and magnetite content was 63.10%, 65.46%, and 63.62%, respectively. The content of iron ore powder 4 is mainly hematite, occupying 46.36%. Iron ore powder 4 has high viscosity. Adding a small amount of iron ore powder can not only replace bentonite, but also improve the quality of raw materials.

Experimental Method 1. The pore volume, pore diameter distribution and average pore diameter of raw materials were measured by nitrogen adsorption instrument. 2. Pelletizing conditions: Using 500 × 150 mm disc pelletizer to implement the pelletizing experiments, linear velocity is 0.98 m/s, and the obliquity is 48°. 3. Determination method for green ball performance. (1) Determination of drop strength of green ball: selecting 20 green balls with diameter of 10.0–12.5 mm; which land freely on the steel plate from a height of 500 mm, then recording the number of non-breaking balls, taking the arithmetic mean of the total number of 20 balls as the index of drop strength (times/balls). (2) Determination of compression strength of the green ball: 20 green balls with a diameter of 10.0–12.5 mm were selected. The compressive strength of each ball was measured on the MTLQ-RQT-1 pellet pressure test machine. The arithmetical average of the compressive strength of 20 balls was taken as the compressive strength index (N/balls). 4. The green ball burst temperature was measured by the measuring device for green ball burst temperature. 5. Roasting process and determination of compressive strength: The experiments of pellet roasting were carried out in vertical electric furnace. The compressive strength of the pellet was determined according to the method for determination of the compressive strength of the iron ore pellets (GB/T14201-1993). 6. Method for determination of metallurgical properties of pellets.

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(1) Determination of reduction degradation property at low temperature: The determination was conducted according to “the method of using cold drum after static reduction of iron ore low-temperature pulverization experiment” (GB/T 13242-91). (2) Determination of pellet reduction property: The determination was conducted according to “method for determination of iron ore reducibility” (GB/T 13241-1991). (3) Pellet reduction expansion: The determination was conducted according to “method for determination of relative free expansion index of iron ore pellets” (GB/t13240-91). (4) Determination of softening properties under load: In the heating process, the temperature at which the material column shrinks by 10% is the softening starting temperature (T10% ), and the temperature at which it shrinks by 40% is the softening ending temperature (T40% ), and the temperature range of softening temperature is calculated by ◿T  T40% −T10% . The sample size was 2–3 mm; the load was 1 kg/cm2 , and the height of the material column was 40 mm. 7. The samples were prepared by Leica prototype machine, and the mineral composition and microstructure of magnesium fluxed pellets were analyzed by Leica DM4500P microscope in Germany.

Experimental Scheme and Results Analysis Test Results of Raw Material Base Performance Granularity Distribution The particle size composition of the raw material was determined by the method of sieving with the square mesh screen and weighing with the electronic scales. The composition of raw material granularity is shown in Table 4. The particles in iron ore powder 1, 2, 3, and 4 which are below 0.074 mm are 31.8%, 83.6%, 77.7%, and 30.5%, respectively. The particle size composition of limestone is mainly between 0.15 and 0.074 mm and those less than 0.074 mm account for 17.2%. The particle size composition of caustic-burned magnesia powder is mainly between 0.15 and 0.045 mm. Adding the proper amount of caustic-burned magnesia powder can improve the particle size composition of the mixture and increase the ball formation rate. But too much will cause the reduction of the burst temperature of the green ball.

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Table 4 Size distribution of raw materials/% Raw material

>0.15 (mm)

0.15–0.074 0.074–0.045 85%). Low resistivity at high growth temperature is supposed to be attributed to the high carrier concentration resulting from its crystallinity and molecular orbital ordering. On the other hand, at relatively low growth temperatures below 150 °C, indium oxide behaves as a transparent semiconducting oxide (TSO) so that the related TFT devices exhibit reasonably high performance with a saturation mobility exceeding 15 cm2 /V s and a threshold voltage near 0 V. It should be noted that not only can the composition of the combined ALD thin film be adjusted by changing the number of sub-cycles, but the influence of the growth mechanism on the composition can be enhanced if the surface termination is changed using different metal precursors. It was reported that a discrepancy appeared between the actual growth of IZO and the sum of the respective binary oxides, which increased at higher deposition temperatures. Two surface reaction cases were proposed for the In2 O3 cycle, where higher temperature may cause the InCA-1 precursor to release an additional CH3 ligand so that In bonds to two −OH sites. During the following ZnO

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Fig. 1 a Schematic illustration of the cross section of a flexible TFT structure incorporating InOx as the active layer. b Representative transfer characteristics (ID vs. VG ) of the devices incorporating InOx films grown at 125 and 150 °C

cycle, the number of −OH sites on the surface is reduced, resulting in a suppressed ZnO growth rate at high growth temperatures despite the increased reactivity [5]. The IGO TFTs were fabricated in a top gate bottom contact structure, as shown in Fig. 2, using different supercycles composed of InOx and GaO sub-cycles to control the atomic composition as well as electrical characteristics, such as carrier concentration and resistivity. The ALD IGO TFT achieved an optimum mobility of 9.45 cm2 /V s with a 3InO–GaO supercycle, which is comparable to that achieved in previous reports of sputter deposited IGO TFTs [6]. Recently, many researches have started to pay attention to the effect of hydrogen in oxide semiconductors, as it tends to form –OH bonds, each generating a free electron; thus, hydrogen acts as a donor. Such a mechanism involves the formation of −O2− H+ charge states, which induce excess free carriers in oxide semiconductors. In addition, hydrogen in the insulator layer can diffuse into the oxide semiconductor active layer during the deposition process or post-treatment process. However, the specific role of hydrogen in the device performance is still not clear. Park’s group compared the concentration of hydrogen in SiO2 thin films grown by thermal oxidation, PECVD and ALD, and found that it differs depending on the deposition method, as shown in Fig. 3a [7]. Thus, to figure out the effect of hydrogen on device performance, TFTs based on sputtered IZTO semiconductors were fabricated using thermal SiO2 , PECVD SiO2 , and PEALD SiO2 as gate dielectrics, where superior device performance and stability were observed in the last case (Fig. 3b). A linear field effect mobility of 68.5 cm2 /(V s) and a net threshold voltage shift (Vth ) of approximately 1.2 V under positive bias stress (PBS) were obtained using PEALD SiO2 as the gate insulator. The authors suggested that the relatively

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Fig. 2 a Structure diagram of top gate bottom IGO TFT. b Transfer performance of the InO–GaO sequence in different InO sub-cycles

Fig. 3 a Elastic recoil detection from SiO2 deposited by thermal oxidation, PEALD (150 °C), and PECVD with an ITZO layer deposited onto them, b ITZO TFT transfer curves with respect to the type of SiO2 gate insulator, c proposed effect of hydrogen in ITZO TFTs based on thermal, PEALD, and PECVD SiO2 gate dielectrics

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Fig. 4 a Variation in device parameters: threshold voltage, field effect mobility, and subthreshold slope as a function of bending cycles for flexible IZO TFTs fabricated on plastic substrates with bending radii of 5 and 2.5 mm. b Illustration of flexible ALD TFT experimental rolling test and device parameter shift under repetitive rolling cycles with bending radius of 1.5 mm

high concentration of hydrogen in the PEALD SiO2 induces a high carrier density in the ITZO layer deposited onto it, which results in enhanced charge transport properties, as shown in Fig. 3c. Also, it is most likely that the hydrogen atoms have passivated the electron traps related to interstitial oxygen defects, thus resulting in improved stability under PBS. The stability of ALD oxide TFTs under mechanical stress is exhibited in Fig. 4a, where the threshold voltage shift of an ALD IZO TFT after 5000 cycles with much a smaller stress radius (2.5 mm) was smaller than that of InOx ALD TFTs. This is supposed to be related to the decreased number of defect states and greater Zn–O bonding strength of IZO TFTs. Further improvement was observed in IZTO TFTs (Fig. 4b), where almost no performance degradation appeared even after 240,000 bending cycles with a much smaller radius (1.5 mm), which is quite superior to sputtered oxide semiconductor TFTs. Although the reason for this difference in flexibility is still not clear, ALD TFTs exhibit great potential for flexible devices, such as rollable, foldable, and stretchable displays.

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Conclusion Recent progress in oxide TFTs has been discussed with a focus on n-type active layers grown by ALD, anticipating the development of device performance. An n-type active layer can be improved by oxide semiconductor combinations from binary (ZnO) to ternary (IZO, IGZO, ITZO) compounds, since the ALD method allows for easy control over the atomic composition of nanostructured layers that can result in decreased bulk active defect states and improved device mobility as well as increased carrier concentration. In addition, some research challenges were discussed in terms of gate insulators and flexible TFTs fabricated using ALD techniques, suggesting that functional films grown via ALD, with their high quality and low deposition temperature, could offer possible solutions for applying oxide semiconductor TFTs for next-generation displays.

References 1. Sheng J, Jeong HJ, Han KL, Hong TH, Park JS (2017) Review of recent advances in flexible oxide semiconductor thin-film transistors. J Inf Disp 18(4):159–172 2. Park J, Mange W-J, Kim H-S, Park JS (2012) Review of recent developments in amorphous oxide semiconductor thin-film transistor devices. Thin Solid Films 520:1679–1693 3. Sheng J, Han K-L, Hong T, Choi W-H, Park JS (2018) Review of recent progresses on flexible oxide semiconductor thin film transistors based on atomic layer deposition processes. J Semicond 39:011008 4. Sheng J, Choi D-W, Lee S-W, Park J, Park JS (2016) Performance modulation of transparent ALD indium oxide films on flexible substrates: transition between metal-like conductor and high performance semiconductor states. J Mater Chem C 4(32):7571–7576 5. Sheng J, Lee H-J, Oh S, Park JS (2016) Flexible and high-performance amorphous Indium Zinc Oxide thin-film transistor using low-temperature atomic layer deposition. ACS Appl Mater Inter 8(49):33821–33828 6. Sheng J, Park E, Shong B, Park JS (2017) Atomic layer deposition of an Indium Gallium Oxide thin film for thin-film transistor applications. ACS Appl Mater Inter 9(28):23934–23940 7. Sheng J, Han J-H, Choi W-H, Park J, Park JS (2017) Performance and stability enhancement of In−Sn−Zn−O TFTs using SiO2 gate dielectrics grown by low temperature atomic layer deposition. ACS Appl Mater Inter 9(49):42928–42934

Adsorption of Fluoride Gases in Aluminum Production by Using of Nanotechnology Mohsen Ameri Siahooei and Kambiz Bordbari

Abstract HF, CF4 , C2 F6 and SiF4 are the main fluoride gases evolved in the Hall –Heroult process. However, the major contributor of fluoride is HF. Since these gases are very toxic, they must be adsorbed from the environment. In this research, adsorption of HF, CF4 , C2 F6 and SiF4 in aluminium smelter is discussed and compared. Due to the capability of nanotubes in gas adsorption, this study has been conducted to figure out the adsorption of fluoride gases on (8,8) armchair carbon nanotubes (CNTs). Lennard-Jones potential was used for gas–gas and gas–carbon nanotube interactions. In addition, the potential parameters for the carbon–gas and carbon— carbon interactions were obtained from the Lorenz-Berthelot combining rules. The simulation results showed that this adsorption can be a possible solution for separation of toxic gases from the environment. The proposed method provides a new horizon in the aluminium industry. Keywords Carbon nanotube · Monte Carlo simulation · Adsorption gas

Introduction Perfluorcarbon (PFC, e.g. tetrafluoromethane and hexafluoroethane) emissions are harmful to the environment because of their global warming potential and, therefore, it is a challenge for the aluminum industry to reduce the evolution of such gases. Tetrafluoromethane, CF4 , and hexafluoroethane, C2 F6 , are greenhouse gases emitted when anode effects occur during primary aluminium production (1). Accurate accounting and inventory of PFC emissions is increasingly important as primary aluminum producers have made commitments, either made voluntarily or M. A. Siahooei (B) Almahdi-South Hormoz Aluminium Smelter, P.O. Box: 79171-7-6385, Bandar Abass, Iran e-mail: [email protected] K. Bordbari Materials Science and Engineering Department, Shahid Bahonar University of Kerman, P.O. Box: 76135-133, Kerman, Iran © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_11

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based on national regulatory requirements, to meet GHG emissions objectives. Tetrafluoromethane (CF4 ) and hexafluoroethane (C2 F6 ), are emitted from aluminum smelters during anode effects (AEs). An anode effect occurs when the current density in the anode exceeds the critical current density. The critical current density depends on temperature and chemical composition of the electrolytic bath. Decreasing alumina content in the electrolytic bath is the most important factor and when it gets too low (below 2%), the cell voltage rises and bath and carbon anodes begin to react. Aluminum smelters are considered to be the largest anthropogenic source of PFC emissions worldwide. In 1991, Iijima announced the discovery of multiwalled carbon nanotubes as a byproduct in fullerene production [1, 2]. Since then, great efforts were made to improve the yield during the preparation and purification. Single-walled carbon nanotubes (SWNT) are available since 1992 in enough quantities to be studied and immediately attracted attention because of potential technological applications and also from a basic research perspective because it was the first from of carbon that could provide physical realizations of ideal systems. Single-walled carbon nanotubes could be considered as the result of bisecting a C60 molecule at the equator and the two resulting hemispheres are joined with a cylindrical tube one monolayer thick and with the same diameter as C60 . If the C60 molecule is bisected normal to a fivefold axis, an armchair tubule is obtained, whereas if the bisection is made normal to a 3-axis, a zigzag type tubule is obtained. Other chiral tubles can be formed with a screw axis along the axis of the tuble. Carbon nanotubes can be specified in terms of the tuble diameter dt , and chiral angle θ , that define the chiral vector. ch  na1 + ma2 . The tube can then be identified using a pair of integers (n, m) that define the chiral vector [3]. Carbon nanotubes have been found to assemble in bundles where the tubes are in a hexagonal array with different lengths. Carbon nanotubes have gathered much attention both from fundamental science and technological interests. Very high chemical stability and mechanical strength made the carbon nanotube a very important material in nanotechnology. Existing theoretical literature suggests that defect-free, pristine carbon nanotubes (CNTs) interact weakly with many gas molecules like H2 O, CO, NH3 , H2 , and so on [4]. In this work, grand canonical Monte Carlo (GCMC) method is used to study the Fluoride compounds adsorption gas on carbon nanotube. Single-walled carbon nanotubes are selected to be the adsorbent. To make a comprehensive work, the influence of temperature as well as pressures on the adsorption is also studied. The simulation results in this work can be used to optimize the Fluoride compounds adsorption at given pressures and temperatures.

Simulation Method The Monte Carlo statistical mechanical simulation were carried out in a standard manner using the Metropolis sampling technique in canonical (T, V, N) ensemble. In this work, all of the particles include sulfur compounds molecules, and carbon atoms

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123

are treated as structureless spheres. Particle–particle interactions between them are modeled with Lennard-Jones potential located at the mass center of the particles. In this work, as in the works of many researchers, the cut and shifted Lennard-Jone (LJ) potential was used to represent the interaction between sulfur compounds molecules.  φff (r) 

φij (r) − φij (rc ) r < rc 0 r ≥ rc

(1)

where r is the interparticle distance, rc is the cut off radius, rc  5σff · φij is the full LJ  12  6  potential, φij  4εff σff /r − σff /r , where εff and δff are the energy and size parameters of the fluid. They are 460 and 2.75 nm for HF also 220 and 5.3 nm for C2 F6 also 134 and 4.662 for CF4 also 171.9 and 4.88 nm for SiF4 here respectively [5–7]. The interaction between the wall and a hydrogen sulfide molecule is calculated by the site-to-site method [8, 9]. Uf w



 Nf Ncabon    σf w 12 σf w 6  4εf w − rij rij i1 j1

(2)

where Nf is the number of HF gas molecules, Ncarbon is the number of carbon atoms of the wall of SWNT. εf w and δf w are the cross-energy and size parameters, which are obtained from the Lorentz-Berthelot (LB) combining rules. Energy and size parameters of carbon atoms are 28.0 and 0.34 nm, respectively [10]. rij is the distance between a gas hydrogen fluoride molecule and an atom of the wall of SWNT. Lorentz-Berthelot rules are used to calculate the parameters of interaction between different kinds of particles. In this calculation, all of the particles are regarded as spheres. Interaction among particles is modeled with Lennard-Jones potential acted on the mass center. The initial configuration was generated randomly (Fig. 1). For a fixed cell, three types of moves were used to generate a Markov chain, including moving, creating, and deleting a molecule and make new configurations. The three types of moves have the same probability and each has different receiving opportunities. Configurations  are accepted when they obey Metropolis, Sampling scheme where E is the change of total energy in the system in proportion to exp −E KT (Fig. 2).

Fig. 1 Initial configuration

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Fig. 2 Armchiar (8,8)

To ensure good thermodynamical averages, for a single isotherm point typically 5 × 106 moves have been performed to equilibrate the system. For each of five hundred configuration, one configuration is selected, and names snapshot. Diagram energy of produced configuration to a number of snapshots show that system reaches to the equilibrium (Fig. 3).

Fig. 3 Snapshot to percent of abundance

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125

Fig. 4 Snapshot for second half to percent of abundance

The ensemble average energy of a system for the second half of snapshots is drawn, and initial part is discarded. Because of initial part far away to the equilibrium. Diagram of energy for the second half, show that, the system reach to equilibrium (Fig. 4). The statistical error has been reported in this work. STDEW is the standard deviation of the calculated average in the simulation of eight number is 0.64% (simulation error). The dimensions of simulation cell are (200 × 100 × 34.5) Å. We considered single-walled armchair (8,8) nanotubes with an open edge (Fig. 2). The number of carbon atom is 320. The diameters of the nanotubes are 10.854 Å, and the average bond length is 24 Å respectively. The number of molecules gas calculated by the virial equation of state and input to the GCMC calculation. The equation of state of real gases is best represented, by the series (Eq. 3)  B(T ) C(T ) D(T ) + + + ··· (3) PVm  RT 1 + Vm Vm2 Vm3 where B(T ), C(T ), and D(T ) are respectively termed the second, third, and fourth virial coefficients. (P) is the pressure, (V m ) is molar volume, (T) the absolute temperature, and (R) the gas constant [11].

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Fig. 5 Distribution of C2 F6 adsorption density versus the distance from the axis of nanotube at different temperatures and pressure of 11 MPa

Results and Discussion Comparative investigation of adsorption rate of CF4 , C 2 F6 , HF, SiF4 gases at the same pressure and different temperatures in carbon nanotubes using Monte Carlo simulation In this study, the effect of temperature on adsorption of fluoride gases compounds of aluminum electrolysis process in single-walled carbon nanotubes is investigated. The first gas is C2 F6 and used temperatures are: 180, 323 and 195.1 K and constant pressure is equal to 11 MPa. Figure 5 indicates the density distribution of C2 F6 gas in terms of distance from the axis of nanotube at different temperatures. The number of gas molecules in this situation is calculated by Second virial coefficient [12] and Lennard-Jones potential parameters are extracted [10]. As it can be seen from Fig. 5, increasing of temperature decreases the density of adsorption on carbon nanotubes. At a temperature of 180 K, maximum distribution of C2 F6 gas is around the nanotube wall. In this case, the maximum temperature and maximum absorption peak of density are calculated. The density decreases with distance from the axis of the nanotube. Table 1 shows the number of gas molecules presented in simulation cells, which is calculated by the second Virial coefficient using the specified temperature. According to Table 1 it is clear that the total density of C2 F6 gas is reduced by increasing temperature. In Fig. 6 the changes of C2 F6 density is plotted versus temperature. The negative slope of the graph in this figure represents the C2 F6 density decreases with increasing temperature. Figure 7 shows the adsorption density of CF4 gas versus the distance from the axis of nanotube at different temperatures. The number of gas molecules is calculated by Second virial coefficient [13] and Lennard-Jones potential parameters have been taken [10]. The applied temperatures are: 225, 323, and 800 K and pressure is constant and equal to 11 MPa. As it can be seen from Fig. 7, with increasing of temperature the adsorption density on carbon nanotube is reduced and most of the CF4 gas distribution is around the nanotube wall at 225 K.

Adsorption of Fluoride Gases in Aluminum Production … Table 1 The inside, outside and total density of C2 F6 adsorption in single-walled carbon nanotube at different temperatures

127

Temperature (k)

180

195.1

323

Second virial coefficient (cm3 /mole)

−815

−611

−199

Number of molecules

3053

2819

1701

Inside density (molecules/nm3 )

2.25

2.24

1.96

4.99

4.68

4.99

4.82

4.53

3.62

Outside density

(molecules/nm3 )

Total density (molecules/nm3 ) Fig. 6 The effect of temperature on C2 F6 gas adsorption

Fig. 7 Distribution of CF4 adsorption density versus the distance from the axis of nanotube at different temperatures and pressure of 11 MPa

Table 2 indicates the number of CF4 gas molecules calculated in specifies temperature. From this table and Fig. 8 descending process of gas adsorption on carbon nanotube is observed with increasing temperature. The third gas under investigation is HF. Figure 9 shows adsorption density of HF versus the distance from the axis of the nanotube at different temperatures. The number of gas molecules is calculated by Second virial coefficient [5] and the Lennard-Jones potential parameters have been taken [10]. Temperatures used are: 275, 320, 440 K, and constant pressure is equal to 7 MPa intended. As it is clear from Fig. 9, increasing the temperature decreases the density of adsorption on carbon nanotubes. At 227 K, the maximum amount of HF distribution is along the axis of nanotubes which is similar to the CF4 and C2 F6 gases.

128 Table 2 The inside, outside and total density of CF4 adsorption in single-walled carbon nanotube at different temperatures

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Temperature (k)

225

323

800

Second virial coefficient (cm3 /mole)

−815

−611

−199

Number of molecules

3053

2819

1701

Inside density (molecules/nm3 )

2.31

2.12

1.29

4.65

3.16

1.08

4.51

3.09

1.09

Outside density

(molecules/nm3 )

Total density (molecules/nm3 )

Fig. 8 The effect of temperature on CF4 gas adsorption Fig. 9 Distribution of HF adsorption density versus the distance from the axis of nanotube at different temperatures and pressure of 7 MPa

Similar to C2 F6 and CF4 gases, the amount of HF gas adsorption decreases with increasing temperature. This can be seen in Table 3 and Fig. 10. We continued our study with SiF4 gas. In Fig. 11 adsorption density of SiF4 versus the distance from the axis of the nanotube is investigated at different temperatures. The number of gas molecules is calculated by Second virial coefficient [14] and the Lennard-Jones potential parameters have been considered [10]. Applied temperatures are: 293, 323 and 450 K and pressure is constant and equal to 7 MPa.

Adsorption of Fluoride Gases in Aluminum Production … Table 3 The inside, outside and total density of HF adsorption in single-walled carbon nanotube at different temperatures

129

Temperature (k)

275

320

440

Second virial coefficient (cm3 /mole)

39.6

12.15

2.07

Number of molecules

1272

1093

795

Inside density (molecules/nm3 )

13.97

4.19

12.69

11.22

6.72

2.23

11.39

7.18

2.89

Outside density

(molecules/nm3 )

Total density (molecules/nm3 )

Fig. 10 The effect of temperature on HF gas adsorption Fig. 11 Distribution of SiF4 adsorption density versus the distance from the axis of nanotube at different temperatures and pressure of 7 MPa

As it can be seen from this figure, increasing the temperature decreases the density of adsorption on carbon nanotubes and the maximum SiF4 distribution along the axis of nanotubes is at 293 K. Despite we applied same pressure and close temperatures for both SiF4 and HF gases, but the inside, outside and total density of SiF4 are much lower than those of HF. Pressure and temperature are relatively similar to those HF. This fact is clear by comparing Tables 3 and 4 and Figs. 10 and 12.

130 Table 4 The inside, outside and total density of SiF4 adsorption in single-walled carbon nanotube at different temperatures

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Temperature (k)

293

323

450

Coefficient second virial (cm3 /mole)

−145.6

−109.4

−34.5

Number of molecules

1193

1082

777

Inside density (molecules/nm3 )

2.19

2.22

2.49

Outside density (molecules/nm3 )

2.94

2.59

1.40

Total density (molecules/nm3 )

2.89

2.56

1.47

Fig. 12 The effect of temperature on SiF4 gas adsorption

Comparative investigation of adsorption rate of CF4 , C 2 F6 , HF and SiF4 at the same temperature and different pressures on carbon nanotubes using Monte Carlo simulation In this section the effect of pressure on the fluoride compounds adsorption in singlewalled carbon nanotubes (armchair(8,8)) in aluminum industry has been studied. We investigated on four gas: C2 F6 , CF4 , HF, and SiF4 at pressures: 3, 13, 7, and 11 MPa, and at a constant temperature of 323 K. Figure 13 shows the distribution of C2 F6 gas in terms of the distance from nanotube axis. The number of gas molecules is calculated with respect to the second factor Virial [12] and Potential parameters are calculated [5]. From this figure, it can be seen that the maximum adsorption near the axis of nanotube has occurred with very close values, but the overall density calculated at a pressure of 13 MPa is almost doubled compared to the pressure of 3 MPa which is due to the adsorption changes in areas away from the nanotube surface. Diagrams in Figs. 14, 15 and 16, respectively, related to gases CF4 , HF and SiF4 . The results indicate the maximum adsorption of the gases occurs in maximum pressure.

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Fig. 13 Distribution of C2 F6 adsorption density versus the distance from the axis of nanotube at different pressures and temperature of 323 K

Fig. 14 Distribution of CF4 adsorption density versus the distance from the axis of nanotube at different pressures and temperature of 323 K

From above figures, it can be found that in addition to the first maximum adsorption peak at the closest point on the surface of the nanotube there is a weak adsorption peak at a short distance away from the first one which is caused by the second regular layer of accumulated molecules. Another notable point is related to HF gas. This gas is very sensitive to increasing pressure for maximum adsorption in the closest point on the surface of the nanotube. Supplemented information in Tables 5, 6, 7 and 8 represent the adsorption density of CF4 , C2 F6 , HF and SiF4 gases at different pressures. According to these tables and Figs. 17, 18, 19 and 20, with increasing gas pressure, the fluoride adsorption on

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Fig. 15 Distribution of HF adsorption density versus the distance from the axis of nanotube at different pressures and temperature of 323 K

Fig. 16 Distribution of SiF4 adsorption density

carbon nanotubes increases and there is a direct relationship between pressure and adsorption density. Comparative study of adsorption of C 2 F6 , HF, SiF4 and CF4 gases in temperature of 323 K and pressure of 7 MPa on carbon nanotubes using Monte Carlo simulation In this section, we applied Monte Carlo simulation for original system. The calculated number of molecules is 1082 in the gas simulation cell. A number of gas molecules is calculated with respect to the second factor of Virial equation [5, 12–14]. Figure 21 shows the density distribution of these gases versus the distance from the axis of the nanotube.

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Table 5 The inside, outside and total density of HF adsorption in single-walled carbon nanotube at 323 K and different pressures Pressures

3

7

11

13

−109.4

−109.4

−109.4

−109.4

Number of molecules

666

1082

1701

2015

Inside density (molecules/nm3 )

2.04

2.22

2.04

2.09

1.38

2.59

2.43

3.89

1.42

2.56

2.35

3.78

Second virial coefficient

Outside density

(cm3 /mole)

(molecules/nm3 )

Total density (molecules/nm3 )

Table 6 The inside, outside and total density of SiF4 adsorption in single-walled carbon nanotube at 323 K and different pressures Pressures

3

7

11

13

Second virial coefficient (cm3 /mole)

12.5

12.5

12.5

12.5

Number of molecules

468

795

1717

2030

Inside density (molecules/nm3 )

2.43

6.19

4.99

7.63

Outside density (molecules/nm3 )

3.00

5.72

5.35

8.20

2.97

5.15

5.15

8.1

Total density

(molecules/nm3 )

Table 7 The inside, outside and total density of CF4 adsorption in single-walled carbon nanotube at 323 K and different pressures Pressures

3

7

11

13

−72.4

−72.4

−72.4

−72.4

Number of molecules

666

1082

1701

2015

Inside density (molecules/nm3 )

2.36

2.67

2.31

2.63

2.18

3.68

4.65

5.00

2.19

3.62

4.51

4.85

Second virial coefficient

Outside density

(cm3 /mole)

(molecules/nm3 )

Total density (molecules/nm3 )

Table 8 The inside, outside and total density of C2 F6 adsorption in single-walled carbon nanotube at 323 K and different pressures Pressures

3

7

11

13

Second virial coefficient (cm3 /mole)

−199

−199

−199

−199

Number of molecules

464

1082

1701

2010

1/91

2.24

1.96

2.16

1.82

2.97

3.62

3.69

1.83

2.92

3.52

3.6

Inside density

(molecules/nm3 )

Outside density (molecules/nm3 ) Total density

(molecules/nm3 )

134

Fig. 17 Pressure effect on C2 F6 adsorption

Fig. 18 Pressure effect on HF adsorption

Fig. 19 Pressure effect on SiF4 adsorption

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Fig. 20 Pressure effect on CF4 adsorption

Fig. 21 Adsorption of fluoride gases at a pressure of 7 MPa and temperature of 323 K

Figure 21 indicates that there is a big difference between the HF adsorption and other gases at the same temperature and pressure which represents that this gas is adsorbed six times more than other gases at the closest point on the inside surface of the nanotube. Figure 22 shows a comparison of the adsorption density of four gases. From this figure it is clear that HF absorption rate is twice of other gases.

Conclusion The effect of temperature and pressure on fluoride gases indicate that gas adsorption in single-wall armchair (8,8) nanotubes will increase with decreasing temperature. Moreover, with increasing pressure, the gas adsorption rate will be added. It was observed that C2 F6 is less sensitive to temperature changes while HF is more sensitive to changes of pressure in compare with other fluoride gases compounds.

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Fig. 22 Comparison of adsorption density of fluoride gases at pressure of 7 MPa and temperature of 323 K

Ultimately, at the same temperature and pressure, HF has the maximum rate of adsorption. In future work, we are going to obtain the practical results of the proposed method.

References 1. Yoo D-H, Rue G-H (2002) Study of nitrogen adsoded on single-walled carbon nanotube bundles. J Phys Chem B 106:3371–3374 2. Nguyen H, Mai V (2006) Effect of NH3 gas on the electrical properties of single-walled carbon nanotube bundles. Sens Actuators B 113:341–346 3. Juan P, Fabian S (2003) N2 physisorption on carbon nanotubes: computer simulation and experimental results. J Phys Chem B 107:8905–8916 4. Andzelm J, Niranjan G (2006) Nanotube-based gas sensors-role of structural defects. Chem Phys Lett 421:58–62 5. Murad S, Mansour KA (1986) A model intermolecular potential for hydrogen fluoride including polarizability. Chem Phys Lett 131(1, 2) 6. Meng L, Duan Y (2006) Site–site potential function and second virial coefficients for linear molecules. Mol Phys 104(18) 7. Johnson CHJ, Spurling TH (1971) Mixture virial coefficients for the Hamann-Lambert model for globular molecules Aust J Chem 24(12):2449–2459 8. Chong G, Gao G-H, Yang-Xin Yu, Mao Z-Q (2001) Int J Hydrogen Energy 26:691–696 9. Simonyan VV, Johnson JK (2002) J Alloy Compd 330–332, 659–665 10. Skelland AHP (1985) Diffusional mass transfer. Wiley, New York, p 482 11. Dymond JH, Smith EB (1980) The virial coefficients of pure gases and mixtures: a critical compilation, introduction 12. Dantzler EM, Knobler CM (1969) J Chem Ithaca 73:1335 13. MacCormack KE, Schneider WG (1951) J Chem Phys 19:845, 849 14. Hamann SD, McManamey WJ, Pearse JF (1953) Trans Faraday Soc 49:351

Experimental Study on Competitive Adsorption of SF6 Decomposed Components on Nitrogen-Doped TiO2 Nanotubes Sensor Jun Zhang, Xiaoxing Zhang, Hao Cui and GuoZhi Zhang

Abstract Gas-insulated switchgear (GIS) using sulfur hexafluoride (SF6 ) as insulating medium has been widely utilized in electrical industry. However, SF6 gas will decompose into several components under partial discharge, like SO2 , SOF2 , and SO2 F2 . TiO2 nanomaterials were reported to have an excellent performance in the gas-sensitive fields. In this paper, we have proposed a new method of online monitoring to detect these products by nitrogen-doped TiO2 nanotubes sensor (N–TiO2 ) in order to easily evaluate the operating condition of GIS. We first set up the experimental platform for this work, and then separately discussed the gas-sensitive response characteristics of N–TiO2 to three by-products. Finally, repeatability of this gas sensor was measured. The results show that the N–TiO2 have a good repeatability as a gas sensor. The sensitivity of the N–TiO2 to SO2 F2 is weakest. Compared to SO2 F2 , the N–TiO2 is more sensitive to both SO2 and SOF2 , and better to SO2 . Keywords Gas-insulated switchgear · SF6 decomposed components · TiO2

Introduction In the long-term operation of GIS, partial discharge would cause SF6 gas decomposing and producing products such as SO2 , SOF2 , SO2 F2 , and HF [1–4]. Since the reaction is irreversible, the products cannot be restored to SF6 gas, resulting in a decrease in insulation performance, which will inevitably be detrimental to the safe and stable operation of the equipment. Therefore, we propose a method of online monitoring, that is, measuring the various decomposition products generated by SF6

J. Zhang · X. Zhang (B) · H. Cui State Key Laboratory of Power Transmission Equipment & System Security and New Technology, Chongqing University, Chongqing 400044, China e-mail: [email protected] X. Zhang · G. Zhang School of Electrical Engineering, Wuhan University, Wuhan 430072, China © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_12

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insulating gas under partial discharge in GIS in real time as a basis for judging the defect condition inside the GIS. The gas sensor method’s significant advantages [5–7] in engineering applications such as portable, lightweight, and implantable electrical equipment make it possible to have a great potential application in online monitoring. However, due to the short development time of the gas sensor method, there are still many problems that should be solved in engineering and practical aspects such as accuracy, sensitivity and stability of product detection. In this paper, we modified the surface of the titanium dioxide gas-sensitive substrate by non-metal nitrogen doping and tested the sensitivity of the modified sensitive material to the characteristic products of the SF6 .

Experimental Methods The gas-sensing test platform includes a gas distributor, a gas-sensing test chamber, an electrochemical analyzer, etc., as shown in Fig. 1. The detection principle of the TiO2 gas-sensing material referred to in this paper is resistance type sensing, that is, the resistance value of sensor that before and after the gas is introduced is measured in real time by an electrochemical analyzer. And then, the computer will complete the work of analyzing the trend of resistance value change and obtain the following important parameters. (1) Resistance change rate (R%) R% can be obtained by measuring the resistance between the two electrodes of the gas sensor. Generally, R% refers to the ratio of the absolute value of the measured resistance value to the initial value at a certain moment. Different types and different concentrations of the gas will cause sensitive changes in R%. Therefore, this feature quantity can provide the most critical information for type discrimination and concentration quantification. The (R%) is usually defined as Computer Gas distributor

Gas duct Electrochemical analyzer

1234

Signal line

Barometer Sensors Gas cylinder

Fig. 1 Gas-sensing test platform

Pump Valve

Gas chamber

Valve

Exhaust gas treatment

Experimental Study on Competitive Adsorption of SF6 Decomposed …

R%  (R − R0 )/R0 × 100%  R/R0 × 100%

139

(1)

where R is the measured resistance value of the sensitive component at a certain moment, and R0 is the initial resistance value of the sensitive component. (2) Response time Response time is the time it takes for the sensor to change from the initial resistance to 90% of the final value. Response time is a key parameter of the sensor. Its value is directly related to whether the sensor can sense the change of measurement quantity in time and respond to environmental changes, thus ensuring that the system prediction can be timely and effective. (3) Recovery time Recovery time refers to the time when the gas-sensitive material reaches the initial value from the changed value. The recovery time reflects the efficiency of desorption of the gas molecules to be tested from the surface of the gas-sensing element, so the shorter the recovery time, the better the desorption performance of the gas-sensing element.

Results and Discussion Response Characteristics to Three Sulfides In this section, the response characteristics of N–TiO2 to three sulfides SO2 , SOF2 , and SO2 F2 were obtained using the experimental platform described in section “Experimental Methods”, using a gas concentration of 50 ppm and a flow rate of 150 mL/min. The gas-sensitive responses of N–TiO2 tested by gas-sensing platform to three sulfides were −65.71%, −29.97%, −1.96%, respectively, as shown in Fig. 2. In order to better understand the change in the results after doping, the results of the intrinsic TiO2 (I–TiO2 ) [8] are also shown in the figure.

Gas-Sensitive Response to Low Concentrations of Three Sulfides Normally, the concentration of three sulfides in the characteristic product of SF6 insulating gas in the actual GIS equipment resulting from partial discharge does not reach the concentration of 50 ppm used in this experiment. The relatively high concentration of the target gas to be tested is selected to obtain more obvious experimental results, which is beneficial to combine the simulation results of the previous work with the gas-sensing characteristics in this paper, and explains the gas-sensing

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Fig. 2 Sensitivity of I–TiO2 and N–TiO2 to SO2 , SOF2 , and SO2 F2

-70

I-TiO2 N-TiO2

-60

R%

-50 -40 -30 -20 -10 0 SOF2

SO2 F 2

Gas Table 1 Gas-sensing parameters of N-doped TiO2 nanotubes to three types of gases at low concentration

SO2

R% (2 ppm)

R% (5 ppm)

R% (10 ppm)

SO2

−3.21

−6.83

−11.54

SOF2

−1.76

−3.51

−7.02

SO2 F2

−0.13

−0.25

−0.40

performance of different TiO2 nanotube-sensitive materials from a microscopic point of view. In view of the fact that the concentration of the gas to be tested in the gas sensor application environment is not high in the practical application of the project, this paper has studied whether N–TiO2 can exhibit a gas-sensitive response to low concentrations of three sulfides, thus laying the foundation for the preparation of TiO2 nanotube sensors with higher measurement accuracy and more practicality in the future. After testing, the response of N–TiO2 to 2, 5, 10 ppm SO2 , SOF2 and SO2 F2 is shown in Table 1.

Repeatability of Gas Sensor Another important performance for sensors is repeatability. In this paper, we have investigated whether the sensor could show equivalent sensitivity to the gas under continuous working conditions. First, the gas to be tested is passed into the detection gas chamber; Second, clean the gas chamber by nitrogen after tested; and finally, repeat the above steps multiple times. This repetitive experiment was carried out with SO2 at a concentration of 50 ppm, and the repeatability curve obtained is shown in Fig. 3. The “*” in the figure indicates that the SO2 gas is introduced at this moment, and “•” indicates that the purge gas nitrogen is supplied at this moment. It can be

Experimental Study on Competitive Adsorption of SF6 Decomposed … Fig. 3 Repeated gas-sensing test of N–TiO2

0

*

*

*

141

12min rest * Ultraviolet * irradiation *

R%

-20

Stop irradiation

-40

-60

0

1000

3000

4000

Time/s

seen from the curve that for the first three times of the introduction of SO2 gas and the cleaning with nitrogen, the resistance value of the N–TiO2 can be restored to the vicinity of the initial value of the resistance. After the sensor is placed for 12 min, the maximum sensitivity of the sensor is less than the response sensitivity of the first three times, which indicate that the sensitivity of the N–TiO2 is not as good as before. This phenomenon may result from a certain degree of nitrogen absorption, which also reminds us of the research, that is, from a practical point of view, whether the sensor will adsorb too much SF6 gas molecules in the SF6 gas atmosphere for a long time and reduce the sensitivity to SF6 decomposition characteristic gas. The sensitivity of the sensor for the fourth time cannot be restored to the vicinity of the first three; it may be that the SO2 gas is chemically adsorbed on the surface of the N–TiO2 . The fifth gas sensitivity test showed that the irradiation of the UV lamp had a positive effect on the desorption of the sensor and returning sensitivity of sensor to normal level. The sixth gas sensitivity test proved this well. In the fifth gas-sensing test, the sensitivity of N–TiO2 continued to decrease during UV lamp irradiation. The reason may be that the high energy provided by the UV lamp makes it difficult for gas molecules to adsorb on the surface of N–TiO2 .

Conclusions In this paper, in order to easily evaluate the operating condition of GIS, we have proposed a new method of online monitoring to detect by-products of SF6 by N–TiO2 . We first set up the experimental platform for this work and then, separately discussed the gas-sensitive response characteristics of N–TiO2 to three by-products. Finally, repeatability of this gas sensor was measured. The results show that the N–TiO2 has a good repeatability as a gas sensor. The sensitivity of the N–TiO2 to SO2 F2 is weakest.

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However, the N–TiO2 is sensitive to both SO2 and SOF2 , which is disadvantageous in the detection of mixed gases because liable to cause cross-sensitivity. Thus, our next work is to improve the selectivity of the TiO2 nanotubes sensor to those three by-products.

References 1. Dong X, Zhang X, Cui H, Zhang J (2017) A first principle simulation of competitive adsorption of SF6 decomposition components on nitrogen-doped anatase TiO2 (101) surface. Appl Surf Sci 422 2. Zhang X, Zhang J, Cui H (2018) Adsorption mechanism of SF 6 decomposition components onto N, F-co-doped TiO 2: a DFT study. J Fluor Chem 3. Zhang X, Zhang J, Dong X, Hao C (2017) A DFT calculation of fluoride-doped TiO2 nanotubes for detecting SF6 decomposition components. SensS 17:1907 4. Beyer C, Jenett H, Klockow D (2000) Influence of reactive SF x gases on electrode surfaces after electrical discharges under SF 6 atmosphere. Dielectr & Electr Insul IEEE Trans On 7:234–240 5. Shendage SS, Patil VL, Vanalakar SA, Patil SP, Harale NS, Bhosale JL et al (2017) Sensitive and selective NO 2 gas sensor based on WO 3 nanoplates. SensS & Actuators B Chem 240:426–433 6. Zhu L, Zeng W (2017) Room-temperature gas sensing of ZnO-based gas sensor: a review. SensS & Actuators B Chem 267 7. Kumar D, Chaturvedi P, Saho P, Jha P, Chouksey A, Lal M et al (2017) Effect of single wall carbon nanotube networks on gas sensor response and detection limit. SensS & Actuators B Chem 240:1134–1140 8. Zhang X, Zhang J, Jia Y, Peng X, Ju T (2012) TiO2 Nanotube array sensor for detecting the SF6 decomposition product SO2. SensS 12:3302–3313

Fabrication of Hardystonite Nano-bioceramic Coating on 306L Stainless Steel Substrate Using Electrophoretic Method and Evaluation of Its Corrosion Resistance to Improve Medical Performance Iman Bagherpour Abstract Metals and alloys are widely used in dentistry, medicine and restoration of defected bone as artificial implants or restorative materials. Bone implants are mainly made of metals to be able to withstand mechanical stresses during operation. Stainless steel 306L is the most common alloy used in the manufacture of bone implants. The main characteristic of this alloy is to have good mechanical properties, but there is always concern about the corrosion resistance of them in physiological and bioactive solutions. In this study, Hardystonite nano-bioceramic was prepared by sol-gel method and after surface preparation was applied on 316L steel at 3 and 5 min time durations and 30 and 50 V voltages using electrophoretic method. The corrosion resistance of 316L stainless steel in ringer solution was measured by potentiodynamic polarization test at constant temperature of 37 °C before and after coating process which due to very good bioactivity of Hardystonite ceramic the corrosion resistance of 316L steel has also increased. The surface microstructure of the coated specimens was investigated using a scanning electron microscope (SEM). The composition of phases and elemental characterization of the coatings was determined by X-ray diffraction method. Results show that the obtained coating on 316L steel is nearly uniform and has no apparent defect at 50 V and 5 min. The Hardystonite coating has improved the corrosion resistance of the substrate so that the corrosion current density in the coated samples is less than the uncoated ones and the corrosion resistance of 316L steel has increased 9 times. The results showed that Hardystonite bioceramic coating applied by electrophoretic method can improve the corrosion behavior and consequently the biocompatibility of metallic implant in medicine applications. Keywords Hardystonite · 316L steel · Electrophoretic · Sol-gel · Ringer solution

I. Bagherpour (B) Department of Materials Science and Engineering, College of Chemical and Metallurgical Engineering Shiraz Branch, Islamic Azad University, 71955 Shiraz, Iran e-mail: [email protected] © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_13

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Introduction Metals and alloys are widely used in dentistry, medicine and restoration of defected bone as artificial implants or restorative materials. Bone implants are mainly made of metals to be able to withstand mechanical stresses during application. Titanium alloys, cobalt based alloys and 306L stainless steel are the most common alloys which are used in the manufacture of bone implants [1]. The main characteristic of these alloys is to have good mechanical properties, but there is always concern about the corrosion resistance of stainless steels in physiological and bioactive solutions. The low bioactivity of these steels means they are not capable of transplanting to living tissue without applying external forces. In contrast to these alloys, there are wellknown ceramics and bioactive glasses that are biocompatible, but they do not have sufficient strength for under-loading applications [2]. In order to provide the high strength required for implants, it is possible to cover metal alloys with biochemical coatings such as calcium phosphate or hydroxyapatite. The ability of such coatings to bone grafting can be useful for the embodiment and fixation of orthopedic prostheses and dental implants [3, 4]. Bioceramic coatings are used to modify the surface of body implants on metallic substrate to increase their corrosion resistance and in some cases are used to create a new surface that provides the implant’s properties and is completely different from the uncoated device [5]. Harystonite (Ca2 Z nSi 2 O7 ) as a calcium silicate-based ceramic can induce biological fixation and tissue ingrowth at the interface of tissue/implant [6, 7]. In addition, Hardystonite in contrast to silicate-based ceramics such as Ca2 SiO4 , Ca MgSi 2 O6 , Ca3 Mg(Si O4 )2 , Ca7 MgSi 4 O16 , Ca2 MgSi 2 O7 are chemically more stable [7, 8]. Hardystonite ceramics can enhance the proliferation of bone marrow stem cells as well as its ability to induce appropriate bone restoration [6]. Hardystonite is chemically more stable than CaSi O3 ceramics and the presence of zinc ions affects ceramic hardness [9]. Hardystonite ceramics support grafting, proliferation and differentiation of human osteoblast (HOB) cells and increase alkaline phosphatase activity, as reported by Ramaswamy et al. [10]. Ducheyne et al. have introduced electrophoretic (EPD) precipitation method to cover hydroxyapatite (HA) on the surface of metal implants and now this technique has attracted considerable attention [11]. The advantages of EPD method include facility, control of the coating thickness, uniformity, low temperature process, economical equipment, medium coating with a complex shape, the ability to form a thick composite film, high purity of sediments, non-occurrence of phase transformation during coating and applicability in medical fields [12, 13].

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Materials and Methods Hardystonite powder was prepared as raw material during the sol-gel process using tetraethyl orthosilicat, (T E O S · Si 4 (C2 H5 O) hexahydrate zinc nitrate(Zn(NO3 )2 . 6H2 O) and tetrahydrate calcium nitrate (Ca(NO3 )2 . 4H2 O). Briefly, it should be mentioned that TEOS was mixed with 2 ml HNO3 and hydrolyzed for 30 min following stirring. Then (Zn(NO3 )2 · 6H2 O) and (Ca(NO3 )2 · 4H2 O) zinc nitrate and calcium nitrate were added to the mixture to obtain the required gel and then held in 60 and 120 °C temperatures for 24 and 48 h, respectively, to obtain a dry gel, then cooked for 3 h in the furnace at 1300 °C to achieve Hardystonite powder. The structure of the Hardystonite phase was evaluated by X-ray diffraction to confirm the crystalline structure. In order to examine the appearance of the produced Hardystonite particles and relative estimation of their size, scanning electron microscopy (SEM) was sued Fig. 2. 316L stainless steel substrate (diameter of 10 mm and thickness of 4 mm) was used for deposition of Hardystonite. The samples were prepared by 60, 100, 220, 400 and 600 sandpapers and then washed for 10 min in an ultrasonic washing machine in order to degreasing and dried at ambient temperature. In order to produce suspension, methanol ((C 2 H 5 OH) merck) solution was used as a solvent. About 70 cc of solvent was poured in beaker and then 3 g of Hardystonite ceramic added to the solvent while stirring. The solution was then placed on agitator at room temperature for 24 h to allow the suspension to reach uniformity. The suspension was placed into the ultrasonic stirring machine for 30 min. An electrophoretic device was used for the coating so that graphite was considered as anode and 316L steel as cathode. During the deposition process, the anode was placed at 1 cm of the cathode and the duration of the coating process was selected 3 and 5 min for 30 V and 50 V voltages, respectively, to obtain a uniform coating. In order to increase the adhesion of the coating to the substrate, the specimens were sintered at 800 °C for 2 h in a furnace under a neutral gas atmosphere of argon. Corrosion behavior of the coated samples was evaluated by Tafel electrochemical test in Ringer solution at 37 ± 1 °C. The samples were used as working electrode and calomel electrode was selected as reference electrode and counter electrode was applied for completing the circuit and performing experiments. In order to adjust the temperature, heating elements and thermometer were used. The potentiostat apparatus was used for testing and the NOVA 1.8 software was used for electrochemical measurements using electrochemical polarization method. After obtaining cathodic and anodic polarization curves for each sample, corrosion potential of the samples was obtained and corrosion current density was determined by Tafel extrapolation method. Then the average corrosion density was calculated for every group of results. Toxicity of Hardystonite bioceramic powder was evaluated by the researchers and non-toxicity of Hardystonite nanoparticles in contact with bone marrow stem cells was confirmed [4, 14].

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Results and Discussion Figure 1 shows the results of Hardystonite phase structure characterization using X-ray diffraction test. It is clear that there are only Hardystonite phase peaks which indicate the formation of a relatively pure and completely crystalline Hardystonite phase. The most notable Hardystonite peaks are located at 2θ  31° of XRD patterns of calcined powders at 1300 °C which shows that the main result has been Hardystonite phase. An elemental analysis of the Hardystonite powder before applying the coating on the 316L steel substrate is shown in (Table 1), as well as in Fig. 2c. These results prove the presence of the main elements (Zn), (Si) and (Ca) in Hardystonite with the desired values in the final composition. Scanning electron microscope (SEM) images in Fig. 2 show that the nanoparticles are formed in very small dimensions, which is consistent with the results of previous studies by the researchers [15]. The size and shape of the bioceramic produced by the sol-gel method are effective in the formation of homogeneous and uniform coatings. The production of Hardystonite nanoparticles in addition to the benefits of homogeneous coatings is also important from other aspects. In many cases, it is necessary to dissolve ceramic powder in other solvents such as silica–sol for the coating process using sol-gel method. In this method, the bioceramic powder in sizes larger than 100 nm is unsolvable in the solution, and even in very small amounts it will not dissolve, which is consistent with the results of other researchers [4]. The structure

Fig. 1 XRD patterns hardystonite ceramics powders prepared in a sol-gel—sintered at 1300 °C for 3h Table 1 Weight percentage of hardystonite constituents measured by pare toxic X-ray diffraction spectrometry

Element

W%

A%

C

8.05

13.42

O

54.12

67.72

Si

11.33

8.08

Ca

13.80

6.89

Zn

12.69

3.89

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Fig. 2 Microstructure of the synthesized bioceramic. SEM picture (a, b) and elemental analysis (EDX) (c) of hardystonite powder

and coating of Hardystonite has been taken at 5000× magnification using scanning electron microscopy (SEM). According to Figs. 3 and 4, it is clear that voltage and time changes have a significant effect on the coating on 316L steel substrate, so that at a constant and different voltage, a voltage of 50 V is the best applied coating and at a constant voltage and variable time, the time of 5 min represent the best and the most coverage in compared to lower voltages and times. The applied Hardystonite coating by electrophoretic method represents a uniform and sound (non-cracked) structure. It can be concluded that electrophoretic method can be applied to metal substrates to obtain coatings with uniform and sound structure. This result is consistent with the results of the other researchers [15]. The achievement of a coherent uniform coating without apparent imperfections is one of the most important goals of the electrophoretic coating process. The presence of such a coating is very necessary to achieve suitable corrosion and biological properties for the substrate.

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Fig. 3 SEM images of Hardystonite coated at a constant voltage of 30 V and variable time on steel 316L. a 30 V and 3 min, b 30 V and 5 min

Fig. 4 SEM images of Hardystonite coated at a constant voltage of 50 V and a variable time on steel 316L. c 50 V and 3 min, d 50 V and 5 min

According to SEM images (Figs. 3, 4, 5 and 6), it can be concluded that at constant voltage, the higher the duration of the test, the more uniform coating with less porosity will be obtained. According to Figs. 3 and 4, it can be concluded that by increasing the voltage, the Hardystonite coating on 316L steel substrate becomes more uniform and without cracks. An elemental analysis of the Hardystonite powder after applying the coating on the 316L steel substrate is shown in (Table 2), as well as in Fig. 7. These results prove the presence of the main elements (Zn), (Si) and (Ca) in Hardystonite with the desired values in the final composition.

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Fig. 5 SEM images of Hardystonite coated, coated in constant time of 3 min and variable voltage on 316L steel. e 30 V and 3 min, f 50 V and 3 min

Fig. 6 SEM images of Hardystonite coated, coated in a constant time of 3 min and variable voltage on steel 316L. g 30 V and 5 min, h 50 V and 5 min Table 2 Weight percent of Hardystonite constituent elements coated on steel 316L, measured by distribution spectrometry

Element

W%

A%

C

4.8

7.76

O

48.67

66.14

Si

13.55

10.49

Ca

21.28

11.55

Zn

12.22

4.06

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Fig. 7 Microstructure of the coated bioceramic. SEM picture (a) and elemental analysis (EDX) (b) and cross-section illustration of hardystonite coated on 316L SS at the sample 50 V and 5 min

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Fig. 8 XRD patterns of hardystonite coated sample at 50 V and −5 min Table 3 The average potential and corrosion density in a ringer solution at 37 °C Samples

βa

βc

Ecorr SCE (V)

St-1

0.12

0.14

−0.14

St-2

0.15

0.15

−0.55

St-3

0.17

0.21

−0.36

St-4

0.14

0.15

−0.37

St-5

0.11

0.14

−0.30

Rp (K cm2 )

Icorr (μA/cm2 )

CR (mpy)

82.52

0.34

0.0038

120.61

0.27

0.0030

185.42

0.22

0.0024

786.08

0.04

0.0004

534.95

0.05

0.0006

In Fig. 8, the results of the characterization of the Hardystonite phase structure are using the X-ray diffraction test. It is clear that there is more Hardystonite Phase that represents the formation of a relatively clean and crystalline Hardystonite phase. The Tafel polarization diagram of 316L stainless steel specimens with and without Hardystonite coating in a ringer solution at 37 ± 1 °C are shown in Fig. 9. The average values of the corrosion current and the corrosion potential of non-coated stainless steel in a ringer solution, as determined by the polarization curves and the Tafel extrapolation method, are presented in (Table 3). According to Table 3 and Fig. 9, it is seen that 316L stainless steel without coating shows a higher corrosion density in the ringer solution (Icorr  0.34). The corrosion density of the coated 316L steel with Hardystonite in the ringer solution has decreased. The results are summarized in (Table 3). Comparison of the curves shows that by applying the Hardystonite coating, the potential leads to more positive values and the corrosion current density decreases. The potential of the coated sample (SCE − 0.37 V) increased in compared to the uncoated sample (SCE −0.14 V). According to Fig. 5 and Table 2, it is noted that the current density of the once coated sample increased by a voltage of 50 V at 5 min in compared to the uncoated substrate.

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-0.1

E vs SCE (V)

-0.2

-0.3

-0.4

-0.5

-0.6

st-1

st-2

st-3

st-4

st-5

-0.7 -9.5

-9

-8.5

-8

-7.5

-7

-6.5

-6

-5.5

-5

log i (A/cm2)

Fig. 9 Toughened Tuff Polarization Test Example St-1  316L without Coating St-2  316L with 30 V Coating in 3 min St-3  316L with 30 V Coating in 3 min St-4  316L with 50 V coverage in 5 min St-5  316L steel with 50 V coverage in 3 min

According to the results reported in (Table 3), it can be concluded that due to the higher corrosion current density of St-1 specimen than the rest of the specimens, the corrosion rate of St-1 has been calculated higher than that of the samples. The corrosion rate of St-4 is also lower than the rest of the specimens. Figure 10 shows the comparison of the corrosion rate of coated and uncoated 316L steel specimens. One of the other parameters that can be extracted from the Tafel curves is the polarization resistance. The value of Rp (polarization resistance) represents the resistance of the system to corrosion. The results show that polarization resistance of St-4 is more than 9.5 times and polarization resistance of St-5 is more than 6 times the polarization resistance of the St-1 sample. Figure 11 shows a better comparison of polarization resistance of the samples. The reason for the increase is the formation of porous coatings having high amounts of pores, which causes pitting and local corrosion of the underlying substrate. Due to increased coating voltage, the amount of ceramic deposited on the substrate increases, and consequently the thickness of the coating is increased. It can be concluded that the corrosion resistance of the substrate increases with increasing the voltage of 50 V and increasing the duration of the coating process.

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0.0045 0.004

St-1

0.003859842

St-2

0.0035

St-3

0.003065169

St-4

0.003

CR (mpy)

0.002497545

St-5

0.0025 0.002 0.0015 0.001

0.000567624 0.000454099

0.0005 0

Fig. 10 Comparison of the corrosion rate in the Hardystonite cover, coated sample, St-1  316L without Coating St-2  316L with 30 V Coating in 3 min St-3  316L with 30 V Coating in 3 min St-4  316L with 50 V coverage in 5 min St-5  316L steel with 50 V coverage in 3 min 900 St-1

786.0811236

800

St-2

Rp (Kohm.cm2)

700

St-3

600

534.9544073

St-4 St-5

500 400 300 185.4248775

200 100

82.52073333

120.6156221

0

Fig. 11 Polarization resistance, Hardystonite resin coated, St-1  316L without Coating St-2  316L with 30 V Coating in 3 min St-3  316L with 30 V Coating in 3 min St-4  316L with 50 V coverage in 5 min St-5  316L steel with 50 V coverage in 3 min

Conclusion Fabrication of Hardystonite nano-bioceramic is possible by sol-gel method. Hardystonite nano-bioceramic can be coated on a 316L stainless steel substrate by an electrophoretic deposition method so that a homogeneous coating without crack can be obtained. Increasing the voltage and coating time will create a thicker and more uniform layer on the 316L steel substrate.

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The optimal conditions for the heat treatment of Hardystonite coating to achieve the best adhesion and strength to the substrate are obtained at a temperature of 800 °C for 2 h under neutral gas atmosphere. Stainless steel 316L coated with Hardystonite ceramic has a higher corrosion resistance than a non-coated substrate. Hardystonite coating had a good effect on the corrosion resistance of 316L stainless steel substrate and reduces the corrosion current density of the substrate. This is equivalent to the increased corrosion resistance of the implant and reduction of its destructive effects on the tissues of the human body.

References 1. Garcia C, Cere S, Duran A (2004) Bioactive coatings prepared by sol-gel on stainless steel 316L. J Non-Cryst Solids 348:218–224 2. Galliano P et al (1998) Sol-gel coatings on 316L steel for clinical applications. J Sol-Gel Sci Technol 13(1–3):723–727 3. Cook SD et al (1988) Hydroxyapatite-coated titanium for orthopedic implant applications. Clin Orthop Relat Res 232:225–243 4. Asgarian R, Doost-Mohammadi A (2016) Evaluation of corrosion behavior, bioactivity and cytotoxicity of nanostructured hardystonite coating on Ti-6Al-4V substrate 5. Hench LL, Wilson J (1993) An introduction to bioceramics, vol 1. World Scientific 6. Zhang M, Lin K, Chang J (2012) Preparation and characterization of Sr–hardystonite (Sr2 ZnSi2 O7 ) for bone repair applications. Mater Sci Eng C 32(2):184–188 7. Gheisari H, Karamian E, Abdellahi M (2015) A novel hydroxyapatite–hardystonite nanocomposite ceramic. Ceram Int 41(4):5967–5975 8. Diba M et al (2014) Magnesium-containing bioactive polycrystalline silicate-based ceramics and glass-ceramics for biomedical applications. Curr Opin Solid State Mater Sci 18(3):147–167 9. Mohammadi H et al (2014) Bioinorganics in bioactive calcium silicate ceramics for bone tissue repair: bioactivity and biological properties. J Ceram Sci Technol 5(1):1–12 10. Ramaswamy Y et al (2008) Biological response of human bone cells to zinc-modified Ca-Sibased ceramics. Acta Biomater 4(5):1487–1497 11. Ducheyne P et al (1986) Structural analysis of hydroxyapatite coatings on titanium. Biomaterials 7(2):97–103 12. Ma J, Wang C, Peng K-W (2003) Electrophoretic deposition of porous hydroxyapatite scaffold. Biomaterials 24(20):3505–3510 13. Zhitomirsky I (1998) Cathodic electrophoretic deposition of diamond particles. Mater Lett 37(1):72–78 14. Wu C, Chang J, Zhai W (2005) A novel hardystonite bioceramic: preparation and characteristics. Ceram Int 31(1):27–31 15. Zhang B et al (2011) Fabrication of nano-structured HA/CNT coatings on Ti6 Al4 V by electrophoretic deposition for biomedical applications. J Nanosci Nanotechnol 11(12):10740–10745

Fabrication of Monodispersed Needle-Sized Hollow Core Polystyrene Microspheres Stanley O. Omorogbe, Esther U. Ikhuoria, Hilary I. Ifijen, Aline Simo, Aireguamen Aigbodion and Malik Maaza

Abstract Among fascinating polymer structures, porous and hollow core structures are usually desirable because of their numerous applications. Herein, we present hollow core structures of polystyrene (PS) microspheres via a simple one-pot novel synthesis using Sodium dodecyl sulfate (SDS) and Cetyltrimethylammonium bromide (CTAB) emulsifiers. The colloidal microspheres are assembled into a non-compact order of colloidal crystal arrays. Microscopic analysis showed that CTAB emulsifier can influence a micelle to act as a soft template for the fabrication of PS particles with inner core hollow interior compared to SDS emulsifier. The optical analysis of the SDS-emulsified PS colloidal crystal films gave wavelengths of 527.011 nm and 642.967 nm that varied from yellow to red colouration as the observation angle changed from 2 to 10 °C, respectively, while CTAB-emulsified PS colloidal crystals showed wavelengths of 556.233 nm and 589.442 nm that corresponds to monochromatic yellow colour that did not change with observation angles. This economic and environment-friendly one-pot procedure can be used to generate a soft-micelle template for the actualization of hollow PS microspheres that may find potential applications as catalysts, fillers, coating materials and drug delivery.

S. O. Omorogbe · H. I. Ifijen · A. Aigbodion Product Development Laboratory, Rubber Research Institute of Nigeria, P. M. B. 1049, Benin City, Nigeria E. U. Ikhuoria (B) Department of Chemistry, University of Benin, Benin City, Nigeria e-mail: [email protected] S. O. Omorogbe · E. U. Ikhuoria · A. Simo · M. Maaza UNESCO-UNISA Africa Chair in Nanosciences/Nanotechnology, College of Graduate Studies, University of South Africa, Muckleneukridge, POBox 392, Pretoria, South Africa S. O. Omorogbe · E. U. Ikhuoria · A. Simo · M. Maaza IThemba LABS-National Research Foundation, Nanosciences African Network (NANOAFNET), 1 Old Faure Road, Somerset West, POBox 722, Somerset West 7129, Western Cape, South Africa H. I. Ifijen Material Science and Technology Division, CSIR-National Institute for Interdisciplinary Science and Technology, Thiruvananthapuram 695019, India © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_14

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Keywords Polystyrene · Monodispersed · Hollow · Sodium dodecyl sulfate · Cetyltrimethylammonium bromide

Introduction Polystyrene (PS) is considered to be a very important polymer in industries due to its rigid, low water absorbing property, low production cost and good processing [1]. This polymer has been used in many industries for catalytic purposes, storage purposes, packaging purposes and as an immobilization agent [2, 3]. However, preparation of PS spheres with fascinating morphologies can be very challenging. With regards to the desirable structure architecture, porous and hollow structures are frequently preferred [4] for polymers because this nature of structure possesses the ability to encapsulate large quantities of guest molecules and potential applications in controlled delivery systems, lightweight fillers, catalysis, coatings, microreactors, etc. [5–8]. Studies have shown that polymer hollow spheres can be realized with the aid of a hard template [9–11]. In this technique, thermolysis or dissolution is used to remove the core template in order to create a hollow interior. The removal of this template can be very challenging. PS hollow microspheres have also been synthesized under γ -ray irradiation at room temperature via irradiation-assisted free-radical polymerizing without the use of any additives [12]. However, if exposed, γ -ray irradiation is very dangerous to our health. Recently, a cheap and environment-friendly one-pot synthesis has been used to form a soft template, such as micelles [13–15] and lipid vesicles [16, 17] for the actualization of polymer hollow spheres. During the synthesis of PS latex, the polydispersity index must be less than 0.5 to ensure good monodispersity as a monodispersed polymeric material has been shown to be more useful than polydispersed latex because of its numerous applications. The level of success in each of these applications is ultimately dependent upon the particle size and its distribution, the morphology of the particles and the surface characteristics among numerous other factors. In order to achieve these factors, monodispersed PS colloidal solutions are usually prepared via emulsion polymerization with or without a surfactant [2]. In this study, the use of SDS and CTAB emulsifiers in facilitating the formation of micelles as a soft template for the fabrication of hollow core interior in PS microspheres was examined.

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Materials and Method Chemicals Styrene, sodium hydroxide, pure nitrogen gas, de-ionized water and potassium persulfate (KPS), sodium dodecyl sulfate (SDS), and cetyltrimethylammonium bromide (CTAB) were all purchased from Sigma-Aldrich and of analytical grade.

Syntheses of PS Microspheres In a typical synthesis, 1.36 g (13.05 mmol) of styrene, 15 g (833.3 mmol) of distilled water, 0.05 g of potassium persulfate (0.185 mmol) and 0.005 g (0.0143 mmol) of cetyltrimethylammonium bromide (CTAB) emulsifier were dispersed in a 50 ml two neck reaction flask and then agitated with a magnetic stirrer (700 rpm). The reaction vessel was maintained at a temperature of 80 °C. At the same time, pure nitrogen was injected into the system. The polymerization process was left for 5 h. The colloidal particles sediment at the bottom after centrifugation and the top layer of clear solvent was poured away. This was repeated for a few times to ensure that the colloidal suspension was free of any unreacted reagent. The procedure for emulsion polymerization of SDS-PS was the same except SDS was used as the emulsifying agent.

Characterization Techniques The morphology of PS microspheres was viewed using a scanning electron microscope (JEOL-JSM 5600 LV), atomic force microscope in the tapping mode (Bruker Multimode, Germany). The functional groups present in the polystyrene (PS) samples were determined using a Perkin-Elmer Series Spectrum Two FT-IR spectrometer over the wave number range of 600–4000 cm−1 . The sample was directly mixed and pelletized with KBr. Powder X-ray diffraction pattern of PS sample was taken using PANalytical EMPYREAN instrument equipped with reference radiation of Cu Kα (λ  1.54 Å) at an operating voltage of 45 kV. The thermal stability of the polystyrene samples was examined with the help of a thermo-gravimetric analyzer, TA Q50, under nitrogen gas atmosphere at a heating rate of 10 °C/min. The examination of the reflectance of colours at 2° and 10° of the colloidal crystals was carried out, using a fiber-optic Maya 2000 Pro spectrometer of Ocean Optics. The optical absorption features of polystyrene colloids in the UV-V is range of 200–800 nm wavelength were measured using a spectrophotometer (SHIMADZU UV-2401PC, Shimadzu, Japan) employing a 1 cm path length quartz cell at room temperature. The average particle size and size distribution (polydispersity index (PDI) were measured

158 Table 1 Effect of surfactant types on the particle sizes of PS latex

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Average particle diameter (nm) Polydispersity index (PDI) Zeta-potential (mV)

PS-SDS

PS-CTAB

479.3

497.5

0.133 −33.80

0.122 −35.80

using Dynamic Light Scattering (DLS) (Nano-Zetasizer, Malvern Instruments) at 25 °C under the scattering angle of 173° at 6333 nm wavelength. The differential scanning calorimetry (DSC) measurement was performed on a DSC 2920 module in conjunction with the TA Instruments 5100 system at a scan rate of 10 C/min under a nitrogen atmosphere.

Results and Discussion Emulsification purpose during polymerization is primarily due to lower surface tension between the organic and aqueous phases, thus, stabilizing the colloidal latex towards the formation of spherically shaped without aggregation [11, 18, 19]. The PS samples were synthesized using SDS and CTAB emulsifiers, while all other reaction parameters were made constant. Table 1 shows the effect of the emulsifiers on the particle sizes, polydispersity, and zeta-potential of the prepared PS microspheres. The average particle diameter of the prepared PS colloidal microspheres is observed to be 479.3 nm and 497.5 nm for PS-SDS, PS-CTAB, respectively as shown in Table 1. The difference in size could be associated with the differences in properties existing between CTAB (cation) and SDS (anion) emulsifier. The dynamic light scattering (Table 1) revealed that a change in surfactant type from SDS to CTAB resulted in a slight increase in the average particle size of the prepared PS samples from 479.3 nm to 497.5 nm. A PDI of // z (building direction) growth. The consistency of ψ-angle at the bottom and top of the sample resulted in similar constitutional undercoolings and evolution of similar grain morphologies.

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Fig. 5 3D APT atom maps showing the distributions of Al, Mg and Si elements inside the melt pools of a bottom and b top specimens. Precipitates are highlighted by iso-concentration surfaces for 3 at.% Mg or 3 at.% Si

Despite the similarities in the microstructural characteristics of bottom and top of DMLS-AlSi10Mg, submicron characteristics of these two locations were different, as shown in Fig. 5. While both Si and MgSi precipitates developed on the top of the sample (Fig. 5b), only MgSi precipitates evolved in the bottom of the sample (Fig. 5a). Moreover, by moving from the top to the bottom of DMLS-AlSi10Mg, fine and dispersed MgSi precipitates are changed to large, needle-shaped ones. The absence of Si precipitates in the bottom of the sample is a result of various heating cycles experienced by the material, which led to dissolution of Si precipitates [12]. On the other hand, MgSi precipitates grew due to diffusion of Mg in the aluminum matrix. It was shown before that Si precipitates play an important role on the strength of DMLS-AlSi10Mg [5, 11]; therefore, their absence can compromise the strength of the alloy. On the other hand, variation of the characteristics of MgSi precipitates can affect the strength of DMLS-AlSi10Mg. Hence, it is essential to investigate the effect of submicron characteristics on the strength of the alloy in detail.

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Conclusions In the current study, the microstructure of a vertically built DMLS-AlSi0 Mg sample was analyzed at two extreme locations; bottom and top of the sample using EBSD and APT techniques to investigate the evolution of micron and submicron characteristics of the alloy during the DMLS process. The microstructure of the sample at the bottom and top was consistent and comprised of columnar grains developed along the building direction. The majority of the grains (76–78%) were columnar in both locations and the grain area was similar. Evolution of {001} fibre texture as a result of // z (building direction) growth was the main reason for columnar structure development, considering the fundamentals of solidification and constitutional undercooling. Despite the similarities in the microstructural characteristics of bottom and top of DMLS-AlSi10Mg, submicron characteristics of these two locations were different. While both Si and MgSi precipitates developed on the top of the sample, only MgSi precipitates evolved in the bottom of DMLS-AlSi10Mg. Moreover, by moving from the top to the bottom of DMLS-AlSi10Mg, fine and dispersed MgSi precipitates are transformed to large, needle-shaped ones. Differences in submicron characteristics can potentially affect the strength of DMLS-AlSi10Mg. Acknowledgements The authors would like to acknowledge Natural Sciences and Engineering Research Council of Canada (NSERC) project number RGPIN-2016-04221 and New Brunswick Innovation Foundation project number (NBIF)-RIF2017-071 for the financial support of this work. The authors would also like to acknowledge Dr. Mark Kozdras at CanmetMATERIALS for facilitating the research. APT analysis was carried out at the Canadian Centre for Electron Microscopy (CCEM), a facility supported by McMaster University, the Canada Foundation for Innovation under the Major Science Initiative program and NSERC.

References 1. Gu DD, Meiners W, Wissenbach K, Poprawe R (2012) Laser additive manufacturing of metallic components: materials, processes and mechanisms. Int Mat Rev 57:133–164 2. Herzog D, Seyda V, Wycisk E, Emmelmann C (2016) Additive manufacturing of metals. Acta Mater 117:371–392 3. DebRoy T, Wei HL, Zuback JS, Mukherjee T, Elmer JW, Milewski JO, Beese AM, WilsonHeid A, De A, Zhang W (2018) Additive manufacturing of metallic components—process, structure and properties. Prog Mater Sci 92:112–224 4. Martin JH, Yahata BD, Hundley JM, Mayer JA, Schaedler TA, Pollock TM (2017) 3D printing of high-strength aluminium alloys. Nature 549:365–369 5. Hadadzadeh A, Baxter C, Shalchi Amirkhiz B, Mohammadi M (2018) Strengthening mechanisms in direct metal laser sintered AlSi10Mg: comparison between virgin and recycled powders. Addit Manuf 23:108–120 6. Olakanmi EO, Cochrane RF, Dalgarno KW (2015) A review on selective laser sintering/melting (SLS/SLM) of aluminium alloy powders: processing, microstructure, and properties. Prog Mater Sci 74:401–477 7. Read N, Wang W, Essa K, Attallah MM (2015) Selective laser melting of AlSi10Mg alloy: process optimisation and mechanical properties development. Mater Des 65:417–424

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8. Basak A, Das S (2016) Epitaxy and microstructure evolution in metal additive manufacturing. Annu Rev Mater Res 46:125–149 9. Kok Y, Tan XP, Wang P, Nai MLS, Loh NH, Liu E, Tor SB (2018) Anisotropy and heterogeneity of microstructure and mechanical properties in metal additive manufacturing: a critical review. Mater Des 139:565–586 10. Hadadzadeh A, Amirkhiz BS, Li J, Mohammadi M (2018) Columnar to equiaxed transition during direct metal laser sintering of AlSi10Mg alloy: effect of building direction. Addit Manuf 23:121–131 11. Hadadzadeh A, Amirkhiz BS, Odeshi A, Mohammadi M (2018) Dynamic loading of direct metal laser sintered AlSi10Mg alloy: strengthening behavior in different building directions. Mater Des 159:201–211 12. Liu YJ, Liu Z, Jiang Y, Wang GW, Yang Y, Zhang LC (2018) Gradient in microstructure and mechanical property of selective laser melted AlSi10Mg. J Alloy Compd 735:1414–1421 13. Persaud SY, Smith JM, Langelier B, Capell B, Wright MD (2018) An atom probe tomography study of Pb-caustic SCC in alloy 800. Corros Sci 140:159–167 14. Yang KV, Shi Y, Palm F, Wu X, Rometsch P (2018) Columnar to equiaxed transition in AlMg(-Sc)-Zr alloys produced by selective laser melting. Scr Mater 145:113–117 15. Gaumann M, Bezencon C, Canalis P, Kurz W (2001) Single-crystal laser deposition of superalloys: processing–microstructure maps. Acta Mater 49:1051–1062

Finite Element Analysis of Particle Pushing During Selective Laser Melting of AlSi10Mg/AlN Composites Marjan Nezafati, Ali Bakhshinejad, Benjamin Church and Pradeep Rohatgi

Abstract Distribution of particles in metal matrix composites has a crucial effect on the efficiency of the mechanical reinforcing process. Re-melting and solidification of the matrix due to laser energy input during selective laser melting increases the probability of the secondary phase agglomeration. Reinforcement particles reallocate locally based on the speed of the solidification front through formation of the melt pool. Adequate energy input and optimized scanning speed are required not only to assure consolidation of the product, but also to control the melt pool geometry and solidification rate and consequently avoid particle pushing and clustering. A finite element model was developed to exhibit the interaction of aluminum nitride particles with the AlSi10Mg melt pool with respect to the solidification front. The model shows that the critical solidification conditions define whether engulfing or particle pushing take place; as an essential consideration when manufacturing metal matrix composites through selective laser melting. Keywords Particle pushing · Selective laser melting · Aluminum · Composite

Introduction Additive manufacturing and specifically, selective laser melting (SLM) is a convenient production method for complicated geometries in applications requiring high dimensional accuracy including metal matrix composite parts for aerospace, turbine blades and biomedical implants [1–3]. The complicated heat, mass, and momentum transfer during SLM process influence the final product quality and require careful study of variables influencing the procedure. Among all different metallic materials, metal matrix composites are precisely vulnerable while being processed through M. Nezafati (B) · B. Church · P. Rohatgi University of Wisconsin Milwaukee, 3200 N Cramer St, Milwaukee, WI, USA e-mail: [email protected] A. Bakhshinejad Medical College of Wisconsin, 8701 W Watertown Plank Rd, Milwaukee, WI, USA © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_31

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SLM method. Melting and re-solidification of the powder via moving laser heat source may push the reinforcement particles causing uneven distribution and decreasing the reinforcement efficiency. There are a considerable number of researches focusing on the behavior of alloys during SLM process, considering the geometry of melt pool, the temperature profile and temperature gradient upon scanning [4–7]. Still, there are only a few studies pointing out the velocity of melt around the impurity particle [8, 9]. In their model, Gu et al. [8] demonstrated that the melt flow velocity vector near the reinforcement particle which is responsible for transportation of the particle is highly sensitive to the SLM processing parameters. However, they focused on the flow velocity rather than solidification velocity. The current manuscript attempts to highlight the importance of solidification rate on the interaction of melt and particle during SLM process applying Finite Element Analysis (FEA).

Model Description A rectangular block with 1 mm × 5 mm × 0.5 mm dimensions in X, Y and Z directions represents the solid base material of AlSi10Mg. A thin layer of powder with thickness of 0.2 mm is added on top of the block. The laser beam heat source is modeled as a heat flux with a Gaussian function distribution scanning the powder layer along the y-axis for the SLM process. The energy input from the heat source or Q is calculated as a function of the heat absorption of the powder, A, the laser power, P, the radius of the laser beam, ω, and the real-time distance of an irradiation point on the surface measured from the laser beam center, r, at time t. Q

  2r 2 2AP ex p − π ω2 ω2

(1)

Considering an ideal spherical geometry for the AlN reinforcement particle, Fig. 1 schematically presents the interaction between the solidification front and the particle based on the model suggested by Rohatgi and Kim [10]. The solidification front advances with a velocity equal to V and the thermal gradient in solid/liquid interface far from the particle is denoted as G. Their analytical model suggests that, if the solidification front moves at velocities higher than the critical value of Vc the particle will be encaptured, otherwise agglomeration of reinforcement particles is expected due to the particle pushing effect. Vc 

σ a0 (k R1 + 1) 18η R1

(2)

Here, σ is the interfacial energy difference, a0 is particle/interface separation behind the particle, k, R1 and η are the curvature of solid/liquid interface behind the

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Fig. 1 Schematic of particle interacting with solidification front

particle, particle radius and liquid viscosity, respectively. The curvature is defined as:   Ga α − 1 (3) k  3 α is the ratio of thermal conductivity of particle to liquid metal. As expressed below,  depends on the solid/liquid surface tension σ SL , melting temperature Tm , density ρ and heat of fusion H. Table 1 reports numerical values of the parameters used in the model. Kp Kl σ SL Tm  ρH α

(4) (5)

Results and Discussion The system of bulk solid and the top powder is preheated to 373 K prior to SLM process to ease the melting of the powder and avoid thermal shocking of the cold base metal and a heat source with 200 W power scans the powder surface in y direction with

342 Table 1 Materials properties and theoretical parameters for particle pushing process

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Parameter

Value

Particle/interface separation behind particle, a0

3 × 10−10 m

Initial temperature, T0

373 K

Radius of laser beam, ω

35 µm

Laser power, P

200 W

Scan speed, V

200 mm/s

Powder layer thickness, tlayer

0.2 mm

Powder porosity

0.4

Solidus temperature, Ts

830 K [11]

Liquidus temperature, Tl

867 K [11]

Latent heat of fusion, H

389 J/g [12]

Density of AlSi10Mg powder, ρ

2650 kg/m3 [6]

Interfacial energy differences, σ

4.56 N/m [13]

Viscosity, η

1.3 × 10−3 Pa s [14]

Laser absorptivity of AlSi10Mg powder, A

0.09 N/m [15]

velocity equal to 200 mm/s. Figure 2 depicts the simulated temperature distribution profile of powder from side, back and top view, respectively. In this model, the region with temperature above the liquidus line or 867 K is considered as the melt pool. As the temperature gradient of the model along the scanning direction shows; the site with laser center focus, demonstrates the highest temperature which is approximately 1400 K. This maximum temperature grantee that the energy input is sufficient for complete melting of the AlSi10Mg alloy and avoid any melting for the AlN particles with melting temperature of 2473 K. As it is illustrated in the side view in Fig. 2, the depth of the melt pool is 0.07 mm, which is not enough to diffuse to the bulk solid and melt the base metal with the current powder thickness. To insure the adequate mechanical strength and metallurgical bonding with the previously fabricated layers, the thickness of the powder layer should be less than 0.07 mm considering all other processing parameters remain constant. Furthermore, back and top views of the melt pool confirm that the width of the molten portion of the powder is equal to 0.11 mm and a hatch distance smaller than this width is required to maintain the continuity in the final product. Figure 2 reports the temperature profile along the scanning path for the frame captured after 20 ms. The temperature profile clearly highlights the mushy zone with the horizontal line between the 3 mm and 4 mm from the starting point, where the temperature range is between 830 and 867 K, which correspond to the solidus and liquidus temperatures, respectively. Due to the small size of the system, the temperature behind the moving heat source remains as high as 540 K which may cause some residual thermal stress in the final product and needs further studies through mechanical analysis and thermal stress calculations in future works.

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Fig. 2 Temperature contour plots during laser scanning at P  200 W and V  200 mm/s from different views a top view b side view c front view d temperature profile on the scanning path along the y-axis

As the scanning heat source moves along the length of the surface, it is essential to track and observe the temperature and phase changes occurring in the system. To fulfill this purpose, the center point on the top surface of X-Y plane is chosen as the representative of the points on the surface that experience the heat absorption from the source. This point is chosen on the topmost surface as it is believed that the highest concentration of energy from the heat source is placed on this plane. Figure 3 shows the temperature variation with time at the center of the topmost layer on X-Y plane during the SLM of AlSi10Mg powder. The sharp peak after 15 ms of laser scanning corresponds to the temperature raise as the laser beam passes over the center point, and the horizontal line following the peak denotes the period the mushy zone is stable after the beam leaves the center point. The constant temperature of 373 K and sharp slope of the plot prior to the laser beam reaching the center point confirm that the conduction demonstrates a minor effect on the temperature profile compared to convection and radiation. Thermal analysis and temperature distribution upon the SLM process provide some essential information required to discuss the melting and solidification of the powder and base metal and consequently, the interaction of the reinforcement particle with the solidifying melt pool. Table 1 summarizes the parameters and materials

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Fig. 3 a temperature profile during the scanning period at the center point of powder surface b geometry of the melt pool at the middle point of the surface powder x  0.5, y  2.5 and z  0.7 mm

Fig. 4 Geometry of the melt pool and solidification front velocity

properties used in this study. The solidification interface has been assumed as the sites with temperature exactly equal to 869 K. The critical interface velocity calculated based on the analytical particle pushing model by Rohatgi and Kim [10] is equal to 5.9 mm/s. As the theory explains a solidification front that advances at a higher velocity than this critical rate, it has the ability to trap and encapture the reinforcement particles suspended in the melt pool in front of the solidification interface. Otherwise with the solidification front velocity slower than this value, the particles are pushed forward to the melt pool and eventually agglomerate at the last solidifying region. The FEA model examines the SLM process to ensure its convenience for processing composite. The FEA model defines the solidification front as a shell with temperature equal to 867 K. With this definition of the solidification front the velocity is calculated for the laser scanning the AlSi10Mg/AlN powder surface to be 193 mm/s. As Fig. 4 shows the solidification front moves at a velocity 30 times faster than the critical interface velocity, which confirms that the AlN particles are engulfed during the laser scanning and no particle pushing will take place.

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Conclusion The temperature profile, melt pool geometry and solidification front velocity are analyzed for the AlSi10Mg/AlN composite system. The maximum temperature in the system is in the range between 1400 and 1600 K which is far lower than the melting temperature of AlN and higher than melting point of AlSi10Mg resulting in suspending the AlN particles in melt pool upon laser melting. The geometry of the melt pool for the studied case suggested that the thickness of powder layer added should be less than 0.05 mm and the hatch distance should be less than 0.11 mm to ensure the continuity of the final product. Comparing the solidification front velocity for SLM process with the critical value calculated for the AlSi10Mg/AlN composite, it was confirmed that the particles will maintain their random distribution during melting and solidification in SLM process as the solidification front advances with a velocity 30 times faster than the critical value and no particle pushing is expected.

References 1. Martin JH et al (2018) Additive manufacturing of metal matrix composites via nanofunctionalization. MRS Commun 1–6 2. Cooper DE et al (2013) Additive layer manufacture of Inconel 625 metal matrix composites, reinforcement material evaluation. J Mater Process Technol 213(12):2191–2200 3. Gu DD et al (2012) Laser additive manufacturing of metallic components: materials, processes and mechanisms. Int Mater Rev 57(3):133–164 4. Dai D, Gu D (2014) Thermal behavior and densification mechanism during selective laser melting of copper matrix composites: simulation and experiments. Mater Des 55:482–491 5. Riedlbauer D et al (2017) Macroscopic simulation and experimental measurement of melt pool characteristics in selective electron beam melting of Ti-6Al-4V. Int J Adv Manuf Technol 88(5-8):1309–1317 6. Li Y, Gu D (2014) Parametric analysis of thermal behavior during selective laser melting additive manufacturing of aluminum alloy powder. Mater Des 63:856–867 7. Yuan P, Gu D (2015) Molten pool behaviour and its physical mechanism during selective laser melting of TiC/AlSi10Mg nanocomposites: simulation and experiments. J Phys D Appl Phys 48(3):035303 8. Gu D et al (2017) A multiscale understanding of the thermodynamic and kinetic mechanisms of laser additive manufacturing. Engineering 3(5):675–684 9. Dai D, Gu D (2016) Influence of thermodynamics within molten pool on migration and distribution state of reinforcement during selective laser melting of AlN/AlSi10Mg composites. Int J Mach Tools Manuf 100:14–24 10. Kim JK, Rohatgi PK (1998) An analytical solution of the critical interface velocity for the encapturing of insoluble particles by a moving solid/liquid interface. Metall Mater Trans A 29(1):351–358 11. Tang M et al (2016) Rapid solidification: selective laser melting of AlSi10Mg. JOM 68(3):960–966 12. Rohatgi PK et al (1994) Evolution of microstructure and local thermal conditions during directional solidification of A356-SiC particle composites. J Mater Sci 29(20):5357–5366 13. Shangguan D, Ahuja S, Stefanescu DM (1992) An analytical model for the interaction between an insoluble particle and an advancing solid/liquid interface. Metall Trans A 23(2):669–680

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14. Brandes EA (1983) Smithell’s metals reference book. Butterworth and Co Ltd., Sevenoaks, Kent, United Kingdom, pp 14.7–14.8 15. Eleftherios L et al (2011) Selective laser melting of aluminium components. J Mater Process Technol 211(2):275–284

Numerical Simulation on the Single-Crystal Grain Structure of GH4169 Superalloy Steel in the Spiral Grain Selector Using Procast Software Zheng Chen, Lan’xin Geng, Yu Yao, Yi Cheng and Jieyu Zhang

Abstract The GH4169 alloy steel is used to fabricate the gas turbine blades. It is required to withstand a gas stream temperature of 1600 °C and have a heat transfer rate of a domestic central heating system. For the regular grain growth under directional solidification condition, we study on how to select a single-crystal grain in the spiral grain selector. The purpose of the numerical simulation technique using the solidification process is not only to obtain the distribution of the temperature field or the flow field, but also to examine the single grain microstructure growing process based on the café module contained in Procast software. The simulation results show that the data basis for the analysis of the solidificationof GH4169 alloy provides details on the grain morphology and a more accurate theoretical result for real as-cast products. Keywords GH4169 alloy · Crystal selector · Solidification structure · Numerical simulation

Introduction GH4169 alloy (IN718 alloy in the USA) has been the most widely used nickelbased superalloy in aircraft engine industry over the past 40 years [1–3]. It has been used in many aircraft engine components, e.g., critical rotating parts, airfoils, supporting structures and pressure vessels, accounting for more than 30% of the total finished component mass of a modern aircraft engine. GH4169 alloy is a precipitation-strengthened iron-nickel-based superalloy that exhibits adequate strength, ductility, and fatigue resistance up to 650 °C. Because of its excellent mechanical Z. Chen (B) · L. Geng Tongling University, Tongling, China e-mail: [email protected] Z. Chen · Y. Yao · Y. Cheng · J. Zhang Shanghai University, Shanghai, China © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_32

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properties and fabricability, GH4169 alloy is a key material in modern aero-engine rotating parts. In the research, the nearly 30-year evolution of materials and processes related to GH4169 alloy was reviewed. These advances are due to, in no small measure, great improvements in both materials and processes that enable improved designs. Indeed, the underlying theme of this article is that processing advances are at least as important as alloy composition development. The spiral selector is the key part for producing single-crystal (SX) blades and ensures the integrity of crystal, which mainly includes starter block and spiral part [4, 5]. In Regel’s work, the influence of spiral part on the grain selection process was studied [6]. Both of the metallographic results and EBSD results proved that the prior location and the special orientation of the second dendrite arms were important for the grains competitive growth during the directional solidification process. Based on the experimental results, two geometrical restrict mechanisms of grain selection were proposed [7]. They were the competitive stimulating effect on the second dendrite arms in horizontal direction, which was resulted from the spiral arc shape, and the growing blocking effect on the primary dendrites in vertical direction, which was resulted from the take-off angle of the spiral part. These models could successfully explain the grain selecting effects of the spiral part. The modified cellular automaton (MCA) technology used to simulate the grains’ competitive growth in spiral part. The changes of grains structure and orientation as the grain growing on were studied. The simulated and experimental results were compared and agreed well. Based on the simulated and experimental results, Influences of structural parameters on the grain selection behavior were proposed. The criteria for designing spiral part were also presented [8–10]. The rapid development of advanced aero-engine and industry gas turbine requires high performance of single-crystal (SX) blade. Spiral selector is very important to produce SX blade, which includes starter block and spiral part. In this research, as the height of grain growth (the distance between studied section and the under surface of the sample) increased, the grain density was changed and the orientation was deviated. Furthermore, the experiment and simulation, and the designing rules for the starter block were given out. Adopting CAFE method, the 3D macro temperature field of solidification process was calculated as well as grain growth [11–13]. The properties of grains competitive growth and evolution process during directional solidification in starter block were analyzed based on macro and micro modeling results, and rules for grains competitive growth was explained.

Simulation Progress The structure of the spiral crystal selector can be divided into three parts: the leading stage, the spiral stage and the transition stage. The structure design of spiral segment is the key to the effect of crystal selection. The main structure parameters of helical segment include helix pitch h s , helix diameter ds , helix angle θ , and helix diameter

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dw. Among them, the pitch h s , diameter ds and spiral angle θ satisfy the relational expression h s  π ds tan θ

(1)

The process of spiral crystal selection consists of two stages. The grain grows at the bottom of the leading-crystal segment randomly, grows and competes under the condition of directional temperature gradient, and finally, only some grains whose optimal orientation is close to the direction of heat flow continue to grow into the spiral segment. After the columnar crystal growth enters the spiral section, it completes the competitive growth under the restricted condition of special structure, and eventually one grain is selected to realize the crystallization process. If two or more grains are selected, the crystallization process fails. In general, into the grain orientation angle of the spiral section is less than 15°. At this point, it is difficult to distinguish different grains by primary arm orientation, so it is necessary to consider the difference of secondary arm orientation to distinguish different grains. As shown in the figure, the tissue experiment results of different cross sections of the spiral section are shown. It can be seen from the figure that when the grain grows into the helix segment, there are more grains. The secondary arms of the grains near the lower edge of the cross section and the inner part of the cross section produced a small portion of lateral growth. As the solidification process proceeds, the secondary arms of the above grains grow further laterally, while the number and size of grains near the upper edge of the helix and the outer edge of the helix decrease in competition, as shown in the figure. With the solidification process, only one grain won the competition and grew into a single crystal. This study for crystal growth in spiral crystal converter is simulated computation, using macro-to-micro coupling algorithm simulation, the macro temperature field grid cell size 1 mm × 1 mm × 1 mm, microstructure calculated with grid cell size 0.1 mm × 0.1 mm × 0.1 mm. The material is the second-generation single-crystal superalloy DS4169 whose chemical composition of the ingot is listed in Table 1, which provided theoretical supports for designing starter block as shown in Fig. 1. As shown in Fig. 1, it is the numerical model of the spiral grain sector and the enclosure simulation results of the grain structure of different cross sections of the spiral section. The results of the thermal simulation during solidification of grain sector in Fig. 2 are basically consistent with the analysis results of the competitive growth process of the crystal geometry constraint growth mechanism shown in Fig. 3. Finally, the grains near the lower edge of the helix pitch h s and the inner part of the helix expand and grow to form a single-crystal structure.

Table 1 The element content of GH4169 superalloy steel Element

Al

Nb

Ti

Cr

Ni

Mo

Fe

Content (weight, %)

0.45

4.82

0.95

17.98

51.69

3.07

~

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Fig. 1 The numerical model of the spiral grain selector

Fig. 2 The thermal simulation result of the refractory alumina, copper (chill) and GH4169

Discussion For the solidification, results was cooling from bottom to the top direction, the thermal gradient curve was calculated in Fig. 4. From that thermal gradient data curve, it could be concluded that the single crystal grew in the condition referred in the paper is compared the model simulation results with the experimental results.

Numerical Simulation on the Single-Crystal Grain Structure …

Fig. 3 The cafe results of GH4169 under directional solidification condition

Fig. 4 The thermal curves on the bottom and top copper position

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Conclusions This article uses the cafe model to simulate the spiral crystal converter the preparation of the second generation of nickel-based single crystal superalloy GH4169. For selected crystal process, the café module method is used to explore the lead crystal section diameter D and crystal section height H and choose crystal section spiral angle α. The aim to compared with the effects of this three parameters is to choose a suitable set for the direct type cylindrical influence. This study aimed to simulate the single-crystal grain structure of GH4169 superalloy steel in the spiral grain selector using the Procast software 2016. The solidification structure under directional cooling condition compared the model simulation results with reported experimental results. By simulating the grain structure in an enclosure, it is found that the simulated café results for the grain structure in the spiral grain sector should be under a well-controlled condition. Acknowledgement Foundation Item: Item Sponsored by Foundation for Natural Science Foundation of Anhui Province Colleges (KJ2016A702), College student maker laboratory construction plan-“Future-oriented”” new material maker laboratory (2016ckjh206) and Tongling college student research fund project (2017tlxydxs080).

References 1. Li H, Fu S, Zhang Q et al (2018) Simulation and experimental investigation of inner-jet electrochemical grinding of Gh4169 alloy. Chin J Aeronaut 31(3):608–616 2. Lu X, Du J, Deng Q et al (2014) Stress rupture properties of Gh4169 superalloy. J Mater Res Technol 3(2):107–113 3. Du J, Lu X, Deng Q et al (2015) Progress in the research and manufacture of Gh4169 alloy. J Iron Steel Res Int 22(8):657–663 4. Sun G, Shang D (2010) Prediction of fatigue lifetime under multiaxial cyclic loading using finite element analysis. Mater Des 31(1):126–133 5. Li P, Cheng L, Yan X et al (2018) A temperature-dependent model for predicting the fracture toughness of superalloys at elevated temperature. Theor Appl Fract Mech 93:311–318 6. Regel G, Putz M, Blau P et al (2016) Influence of microstructures on tribological systems—development of process and surface structure. Proc Cirp 46:281–284 7. Lv P, Sun X, Cai J et al (2017) Microstructure and high temperature oxidation resistance of nickel based alloy Gh4169 irradiated by high current pulsed electron beam. Surf Coat Technol 309:401–409 8. Zhu Z, Zhang LW, Gu SD (2013) Experimental investigation of transient heat transfer between Hastelloy C-276/narrow air gap/silicon steel. Exp Therm Fluid Sci 45:221–226 9. Seli H, Ismail AIM, Rachman E et al (2010) Mechanical evaluation and thermal modelling of friction welding of mild steel and aluminium. J Mater Process Technol 210(9):1209–1216 10. Hu D, Mao J, Song J et al (2016) Experimental investigation of grain size effect on fatigue crack growth rate in turbine disc superalloy Gh4169 under different temperatures. Mater Sci Eng A 669:318–331 11. Li X, Zhang J, Wang B et al (2011) Simulation of stray grain formation during unidirectional solidification of In738lc superalloy. J Cent S Univ Technol (Engl Ed) 18(1):23–28

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12. Guo W, You G, Yuan G et al (2017) Microstructure and mechanical properties of dissimilar inertia friction welding of 7A04 aluminum alloy to AZ31 magnesium alloy. J Alloy Compd 695:3267–3277 13. Xu D, Lu B, Cao T et al (2016) Enhancement of process capabilities in electrically-assisted double sided incremental forming. Mater Des 92:268–280

Powder Packing Density and Its Impact on SLM-Based Additive Manufacturing Taher Abu-Lebdeh, Ransford Damptey, Vincent Lamberti and Sameer Hamoush

Abstract Packing density of metal powders is an important aspect of additive as it directly impacts the physical and mechanical properties of printed products. In order to achieve the most efficient packing of a powder, different grades of that powder must be mixed together in such a way that we minimize the voids. Research has shown that packing the coarser grains first not only yields higher density powders but also decreases balling defects in the finished printed product. In this study, we developed a simple model that adequately predicts the volumetric fractions of different powder grades that can yield the highest powder density. The model accounts for the disparities between theoretical assumptions and experimental outcome, such as volume reduction. The model equations, based solely on void ratios and true specific gravity, will be validated experimentally and compared to other modeling efforts in literature to further prove the potency of the model. Keywords Packing density · Metal powders · Additive manufacturing · Reduction factor

Introduction Efforts to improve the methods of powder metallurgy have been in progress for centuries. In recent years, a heightened interest in metal powder-based additive manufacturing (AM) has emerged and a significant number of researches have been shifted in this direction. Some of these researches investigated powder size distribution as it directly affects the packing density, powder quality, as well as additive manufacturing capabilities. As known, the powder used in additive manufacturing can be produced via several processes such as gas atomization [1, 2], water atomization, T. Abu-Lebdeh (B) · R. Damptey · S. Hamoush North Carolina A&T State University, Greensboro, NC 27411, USA e-mail: [email protected] V. Lamberti Y-12 National Security Complex, Oak Ridge, TN 37830, USA © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_33

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plasma atomization, and plasma rotating among other methods. Each method produces a unique particle size range, and the selection is based on preference and cost option. However, to ensure the most efficient packing, the best particle size distribution that eliminates the most voids in the final product must be chosen. The importance of powder packing is generally appreciated because of its influence on the physical and mechanical properties of the 3D-printed parts, on shrinkage and density on sintering [3], and on the microstructure of the finished product, which may also contribute to the surface finish. Additionally, understanding powder bed properties and its interaction with laser is critical in predicting final part properties. Particle parameters such as powder density, particle size distribution, and shape play an important role in powder bed formation [4, 5]. In order to achieve the densest packing for a powder, different sizes of powder may be mixed together. In the mixing process, the smaller particles will fill in between the interstices of the larger particles, reducing the amount of voids and consequently improving the packing of the powder. Factors that may affect the packing density include loosening effect, wall effect, and wedging effect. In a coarse powderdominant mixture, fine particles are not small enough to fit into interstitial spaces, which will cause loosening effect. When fine powder makes the bulk of the mixture, coarse particles may displace fine particles, but when they are not large enough to fill the spaces, they create voids causing wall effect. Also, when coarse particles are dominant, the wedging effect happens when a fine particle decreases the coordination number of the coarse particles by being in between two coarse particles, rather than in the interstitial space. When fine particles are dominant, the wedging effect occurs when coarse materials impede a layer of fine particle arrangement, creating voids too small for fine particles to fill. Other factors that may affect the packing density include shape and size of the particles, particle size distribution, and powder flowability. Many models have been formulated and studied by researchers in attempts to predict how packing density can be improved. Particle packing models may be categorized into (1) discrete models which can be further categorized based on how many different classes of particles are involved in the mix. Discrete models are classified as binary mixture models such as Furnas model [6], ternary mixture models such as Goltermann or modified Toufar model [7], and multi-component mixture models; (2) continuous models; (3) computer simulation models; (4) design of experiment models; and (5) statistical modeling such as Monte Carlo computation. This research project falls under the category of discrete ternary mixture models.

Project Description Assume, hypothetically, that there are several particle sizes in a system, where the packing of a larger size leaves a certain fraction of voids, and that each subsequent smaller size exactly fills the voids of the preceding size class, without increasing the overall volume of the system. Therefore, if infinitely small particles are successively

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introduced into the system, the packing density could theoretically reach 100%. Now, the question is how many combinations of different sizes are needed to yield the highest packing density? For example, √ in a unimodal (for one size of sphere), the maximum theoretical density is π/(3 2) ≈ 0.74. In a bimodal (two different sizes of spheres), the theoretical highest density could reach 0.933. Thus, hypothetically, adding a third size of smaller diameter than the first two would increase the theoretical packing density (greater than 0.933), and so on. Therefore, introducing an infinite number of decreasing sphere sizes to the first matrix could result in a packing density of 1.0. According to scholars that have worked in this area, a three-component mixture is ideal to achieve the maximum density of powder. Thus, the goal of this project is to develop a simple model that truly represents a three-sized powder packing density, and to generate a simple formula for the volumetric fraction of each particle class required to yield highest density of a mixture. In the proposed model, the packing density of a powder sample is defined as a ratio of the solid volume to the total volume of the sample: Packing densit y(φ) 

Solid volume T otal volume



Vs Vs  1−e Vt Vs + Vv

(1)

where Vs  Volume of solids, Vv  Volume of voids, Vt  Total Volume  volume of solids plus volume of voids, and e  Void ratio (porosity).

Methodology As aforementioned, the highest packing density of a mono-sized sphere in an unbound π ≈ 0.7404, which means the solid material fills only space cannot be more than 3√ 2 74% of the volume space. Consider a container of unit volume with mono-sized spheres (s1 ) crushed into powder such that it the powder occupies approximately 74% of the volume, and the air above is the remainder of the volume. Let the solid π , so that the volume of voids is (1 − a) as shown in Fig. 1. volume be a  3√ 2 Assume we fill up the volume space V 2 with another sphere (s2 ) of smaller diameter such that it reaches the maximum capacity. Based on the theorem from above, π . To filling this column to full capacity would mean a packing density of a  3√ 2 find the packing density () of the mixture, we must find the new solid volume: Vs2  a, which For s1 , 1  VVs1  V1s1  Vs1  a; and for s2 , 2  VVs22  1−a implies Vs2  a(1 − a). So, the total solid volume, Vs  Vs1 + Vs2  a + a(1 − a)  2a − a 2  0.933.

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Fig. 1 Unit volume container filled with powder

(1-a)

V2

V=1 a

V1

This means that the use of two different sizes of spheres increases the theoretical highest density from 0.74 to 0.933. Thus hypothetically, adding a third size of smaller diameter than the first two would increase the theoretical packing density (greater than 0.933), and so on. Therefore, introducing an infinite number of decreasing sphere sizes to the first matrix could result in a packing density of 1.0. It should be noted that the values 0.74 and 0.933 obtained for single sphere packing and the packing of two spheres, respectively, are considered as upper bounds, implying that the actual densities that can be achieved in practice are much lower. Several factors contribute to this difference. These factors are shape, size, and random packing of spheres. Thus theoretically, we can keep adding smaller particles to fill in all the voids. But in practice, this is not possible due to the above-stated reasons.

Packing Density for Ternary Powder Mixtures The packing mechanism in the proposed model is that the smaller particles are introduced into the interstices of the larger packed particles, reducing the voids in the system. Now, assume a stack of powder samples with decreasing sizes (average particle diameter) in a vertical container. Let ri be the volume fraction of the coarser material of any two consecutive powder samples. ri 

Vsi Vsi + Vs(i+1)

(2)

The void ratio e of component i is defined as the ratio of the volume of voids (Vvi ) to the bulk volume (Vi which is the volume of the space that sample i occupies) (Fig. 2): In a unit volume system, the solid volume for component 1: Vs1  1 − Vv1  1 − e1 Similarly, e2 

Vv2 V2



Vv2 Vv1



Vv2 , e1

which implies Vv2  e1 e2

(3)

Powder Packing Density and Its Impact on SLM-Based Additive … Fig. 2 Unit volume container showing solid (powder) volume and void volume

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Vv V1=1 Vs

The solid volume for component 2, Vs2  Vv1 − Vv2  e1 − e1 e2  e1 (1 − e2 ) e3 

(4)

Vv3 Vv3 Vv3   , and thus Vv3  e1 e2 e3 V3 Vv2 e1 e2

The solid volume for component 3, Vs3  Vv2 − Vv3  e1 e2 − e1 e2 e3  e1 e2 (1 − e3 )

(5)

Consider a binary mix of maximum density. Let ri be the volume fraction of the larger of two consecutive particle sizes, di . Thus, the fraction of the smaller of two consecutive particle sizes, di+1 , that will exactly fill the voids of the larger component is 1 − ri . Suppose another set of smaller particles of size di+2 can be introduced into the interstices of the larger particles of size di+1 , and then the total volume of each component size will be given by a series of terms of decreasing magnitude. The first term is ri and the second term is 1− ri .The common ratio i . We call the larger between two consecutive components would therefore be 1−r ri of any two consecutive components as the coarse component, and hence the volume fractions would be termed the coarse volume fraction. We define a geometric series for the coarse volume fractions as      1 − r2 1 − r2 1 − r3 , (1 − r1 ) ,... (6) r1 , 1 − r1 , (1 − r1 ) r2 r2 r3 Similarly, for any two consecutive particle sizes (components), we can define the weight fraction of the coarser material as      1 − ω2 1 − ω2 1 − ω3 ω1 , 1 − ω1 , (1 − ω1 ) , (1 − ω1 ) ,... (7) ω2 ω2 ω3 The weight fraction of component i can be defined in terms of the true specific gravity and the solid volume of component i as

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w1  G 1 vs1  (1 − e1 )G 1 w2  G 2 vs2  e1 (1 − e2 )G 2 w3  G 3 vs3  e1 e2 (1 − e3 )G 3

(8)

Thus, the percent composition of the larger particles that are filled with smaller sizes is given by ω1 

(1 − e1 )G 1 (1 − e2 )G 2 and ω2  (1 − e1 )G 1 + e1 (1 − e2 )G 2 (1 − e2 )G 2 + e2 (1 − e3 )G 3

(9)

The apparent specific gravity G a  (1 − e)G, where G is the true specific gravity. Thus, G a1  (1 − e1 )G 1 and G a2  (1 − e2 )G 2 . So ω1 1   v1 (10) G a1 (1 − e1 )G 1 + e1 (1 − e2 )G 2 1 − ω1 1 ω1 1 (1 − e1 )G 1  −  − G a2 G a2 G a2 G a2 [(1 − e1 )G 1 + e1 (1 − e2 )G 2 ][(1 − e2 )G 2 ] e1  v2 (11)  (1 − e1 )G 1 + e1 (1 − e2 )G 2   2 , simplifies to The third term of the weight fraction series, (1 − ω1 ) 1−ω ω2   2 (1 − ω1 ) 1−ω ω2 e1 e2   v3 (12) G a3 (1 − e1 )G 1 + e1 (1 − e2 )G 2 Calculating the bulk volume for a three-component mixture, 1 e1 + (1 − e1 )G 1 + e1 (1 − e2 )G 2 (1 − e1 )G 1 + e1 (1 − e2 )G 2 e1 e2 1 + e1 + e1 e2 + , thus V  (13) (1 − e1 )G 1 + e1 (1 − e2 )G 2 (1 − e1 )G 1 + e1 (1 − e2 )G 2

V  v1 + v2 + v3 

Equation (13) represents the total volume of a system that is made up of three different sizes and that are each stacked upon another as separate layers. If the particles of the first component, d1 , are relatively large, and all the other sizes (i.e. d2 , d3 , . . .) are very small (ideally, infinitely small), then the voids created as a result of the d1 -sized particles will be exactly filled, and the final volume of the system will be the volume of the first component (largest sized particles). Therefore, if each smaller sized particle exactly fills the voids created by the previously packed larger particle component, the total volume will just be v1 . The reduction in volume as a result of filling the voids of component 1 with components 2 and 3 is equivalent to (V − v1 ). This will be the ideal case. However, for an actual system, the decrease in volume will be less than the ideal case, as components 2 and 3 may not exactly fill the voids created by component 1. This decrease in volume can be expressed as

Powder Packing Density and Its Impact on SLM-Based Additive …

 kd (V − v1 )  kd

e1 + e1 e2 (1 − e1 )G 1 + e1 (1 − e2 )G 2

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 (14)

kd is a reduction factor ranging from 0 and 1. It can only be determined experimentally. Factors that affect the value of kd include size ratio, particle shape, loosening effect, wedging effect, etc. If the small particles are infinitely small (i.e. di+1  di ) that they exactly fill the voids, then kd will be unity. Also, if all the particles are of the same size as the first component, then kd will be zero. The determination of kd will be discussed in detail later. The total volume of the mixture then becomes V f  V − kd (V − v1 ) 

1 + e1 (1 + e2 )(1 − kd ) (1 − e1 )G 1 + e1 (1 − e2 )G 2

(15)

The total weight: 

1 − ω2 w  ω1 + 1 − ω1 + (1 − ω1 ) ω2





1 − ω2  1 + (1 − ω1 ) ω2

 (16)

Substituting Eqs. 8–12 into Eq. 16, the total weight can be expressed as   1 − ω2 (1 − e1 )G 1 + e1 (1 − e2 )G 2 + e1 e2 (1 − e3 )G 3 w  1 + (1 − ω1 )  ω2 (1 − e1 )G 1 + e1 (1 − e2 )G 2 (17) Density 

w (1 − e1 )G 1 + e1 (1 − e2 )G 2 + e1 e2 (1 − e3 )G 3 (1 − e1 )G 1 + e1 (1 − e2 )G 2  ∗ Vf 1 + e1 (1 − kd )(1 + e2 ) (1 − e1 )G 1 + e1 (1 − e2 )G 2

(18) After simplifying, we get densit y 

(1 − e1 )G 1 + e1 (1 − e2 )G 2 + e1 e2 (1 − e3 )G 3 1 + e1 (1 + e2 )(1 − kd )

Percent Composition Let di be the average particle diameter of component i in the mixture composition. For a three-component mixture, we define d1  Coarse material; d2  Medium/intermediate material; and d3  Fine material. From the above-derived formula for V f , we can calculate the volume of each component, and consequently the proportion by volume of each component needed in the mixture to achieve the maximum density of the system: 1 (1 − e1 )G 1 + e1 (1 − e2 )G 2 e1 (1 − kd ) Volume for component of diameter d2  (1 − e1 )G 1 + e1 (1 − e2 )G 2

Volume for component of diameter d1 

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Volume for component of diameter d3 

e1 e2 (1 − kd ) (1 − e1 )G 1 + e1 (1 − e2 )G 2

(19)

The packing efficiency (φ) which is the maximum possible packing degree can be expressed as φ

Vs 1 − e1 + e1 (1 − kd )(1 − e2 e3 ) 1 − e1 + e1 (1 − e2 )(1 − kd ) + e1 e2 (1 − e3 )(1 − kd )   V 1 + e1 (1 + e2 )(1 − kd ) 1 + e1 (1 + e2 )(1 − kd )

(20)

Three Conditions for Density Calculation: (1) Equal void ratios: i f e1  e2  e3  e, then Volume fractions and total volume: v1 

1 e(1 − kd ) e2 (1 − kd ) ; v2  ; v3  (1 − e)(G 1 + eG 2 ) (1 − e)(G 1 + eG 2 ) (1 − e)(G 1 + eG 2 )

Packing efficiency: (2) Same true specific gravity: i f G 1  G 2  G 3  G, then

Volume fractions and total volume: v1 

1 1 e1 (1 − kd ) e1 e2 (1 − kd )  ; v2  ; v3  (1 − e1 + e1 − e1 e2 )G (1 − e1 e2 )G (1 − e1 e2 )G (1 − e1 e2 )G

The packing efficiency: (3) Equal void ratios and equal specific gravities:i f e1  e2  e3  e; andG 1  G 2  G 3  G, then

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Volume fractions and total volume: 1 e(1 − kd ) e2 (1 − kd ) ; v2   ; v3   ; and v1   1 − e2 G 1 − e2 G 1 − e2 G

Packing efficiency:

Determination of kd As aforementioned, k d can only be determined experimentally. To the best of our knowledge, there are no or little data on metal powder to determine the value of k d . However, Furnas [6] conducted an experiment on angular materials using different size ratios of different mixtures, and in each case determined the corresponding contraction in volume that resulted. With this data (shown in Fig. 3), we can generate a quadratic fitting and generated an equation for the parameter kd: kd  1 − 2.62k + 1.62k 2

(21)

Kd

where k is the ratio between consecutive sizes. Relation between Volume Reduction Factor and size ratio

1 0.9 0.8 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0 0

0.2

0.4

0.6

k = ratio between consecutive sizes Fig. 3 Relationship between contraction in volume and size ratio

0.8

1

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Model Verification We tested the proposed model with Furnas model [6] for a three-component mixture. The paper cited above includes an example problem where limestone, fine sand and cement are mixed together. The Furnas model is used to calculate the maximum density, along with the corresponding volume and volume fraction of each component. Void ratios (ei ), specific gravities (Gi ) and average diameters of the components (d i ) were given as follows (Table 1): For this example, Furnas [6] does not consider the volume contraction/reduction parameter, thus in demonstrating the utility of his model, and as such, for comparison purposes, our kd value will be kept as 0. (1 − e1 )G 1 + e1 (1 − e2 )G 2 + e1 e2 (1 − e3 )G 3 1 + e1 (1 + e2 )(1 − kd ) (1 − 0.48)(2.5) + (0.48)(1 − 0.42)(2.65) + (0.48)(0.42)(1 − 0.52)(3.10)   1.39 1 + (0.48)(1 + 0.42)(1 − 0)

densit y 

Volume for limestone, v1   Volume for fine sand, v2   Volume for cement, v3  

1 (1 − e1 )G 1 + e1 (1 − e2 )G 2 1  0.491 (1 − 0.48)(2.5) + 0.48(1 − 0.42)(2.65) e1 (1 − kd ) (1 − e1 )G 1 + e1 (1 − e2 )G 2 0.48(1 − 0)  0.236 (1 − 0.48)(2.5) + 0.48(1 − 0.42)(2.65) e1 e2 (1 − kd ) (1 − e1 )G 1 + e1 (1 − e2 )G 2 (0.48)(0.42)(1 − 0)  0.0989 (1 − 0.48)(2.5) + 0.48(1 − 0.42)(2.65)

These volumes computed with the proposed model are quite closely related to the values obtained with Furnas’ model [6], which, respectively, are 1.39, 0.493, 0.234 and 0.0986. The proposed model gives a total volume of 0.825 (V f ), compared with 0.8260 for the Furnas model. The proportion by volume of each component required to produce the maximum density of the system is compared as follows (Table 2): The values shown above reveal that the proposed model can predict the proportions by volume required to produce concrete of maximum density. Not only does the

Table 1 Data used in model verification Limestone

Fine sand

Cement

d i (in)

2

0.045

0.001

ei

0.48

0.42

0.52

Gi

2.5

2.65

3.10

Powder Packing Density and Its Impact on SLM-Based Additive … Table 2 Comparison between the proposed model and Furnas model [6]

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Proportions for max. density (%)

Proportions –> furnas model (%)

% error (%)

Limestone

v1/Vf  59.5

v1/Vf  59.7

0.34

Fine sand

v2/Vf  28.5

v2/Vf  28.3

0.71

Cement

v1/Vf  12.0

v1/Vf  12

0.00

proposed model accurately calculates the proportions of concrete compositions, but it can also be used in several other applications where different powder proportions are required to create a maximum density mixture.

Conclusions The proposed model proves to be very powerful in its versatility of application. The derivations shown suggest that for any powder sample, we can determine the maximum packing density of powder to be used for any application, notably 3-D printing. This flexibility of the model is very valuable for industrial purposes. Even though research suggests that a ternary mixture is the best to achieve maximum density of a given powder type, the model can be derived for any powder mixture with multiple number of components. Not only can the model be applied to powder mixtures of the same material but this model can be applied to mixtures with different components, as with the concrete example problem that was used for model verification. The following conclusions can be drawn: 1. The size composition of a powder mixture is a very important factor in determining the spreadability of the powder mixture and the density of the 3D-printed product. 2. To attain the most efficient packing, different sizes of the powder must be mixed together in such a way that it minimizes the voids, by allowing the smaller particles fill in between the interstices of the larger particles. 3. A model that predicts the volume fraction of each powder size becomes necessary to predict the maximum possible powder density. 4. Researchers have studied the problem of packing spheres into containers over the years, but to the best of our knowledge there is little success that has been achieved in the mathematical development of the laws of powder packing density. 5. In this research project, we have developed a simple model that adequately predicts the volumetric fractions of powder mixture that can yield the highest powder density. The model accounts for the disparities between theoretical assumptions and practical outcome, such as volume reduction. The model equations are based solely on void ratios and true specific gravity.

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6. The proposed model is a discrete model which assumes that each class of particle will pack to its maximum density in the volume available. 7. Although the proposed model is used to model ternary mixtures, it can be extended to include multicomponent mixtures. 8. For generalization purposes, the model assumes that the particles being packed are perfect spheres, and that the size ratio between successive sizes is constant for the entire system. 9. According to Furnas [6], for a maximum density of three-component sizes, the size ratio must be very small. 10. The proposed model accurately predicted the proportions of concrete compositions for maximum packing density. It can also be used in several other applications where different powder proportions are required to create a maximum density mixture. 11. Although the model is derived for metal powder mixtures, it can be applied to mixtures of other systems such as asphalt concrete, paint, rubber, coal storage, etc.

Recommendations and Future Work 1. The proposed model was derived with the aid of some assumptions; future developments of this model should focus on generating experimental results to further validate it and verify its practical applications. 2. The volume reduction factor k d can be evaluated only experimentally. Its empirical equation was obtained by Furnas limited data on granular material [6]. Experimentation on metal powder may show a more refined and accurate relation. 3. Experimental work should include factors that directly impact the packing density such as shape and size of the particles, particle size distribution (size ratio), wall and loosening effects, “Fine grain dominant,” and “coarse grain dominant” mixtures.

References 1. Pérez-de León G, Lamberti VE, Seals RD, Abu-Lebdeh TM, Hamoush SA (2016) Gas atomization of molten metal: part i. numerical modeling conception. Am J Eng Appl Sci 9(2):303–322 2. Abu-Lebdeh TM, Pérez-deLeón G, Hamoush SA, Seals RD, Lamberti VE (2016) Gas atomization of molten metal: part II. Applications. Am J Eng Appl Sci 9(2):334–349 3. German RM (1992) Prediction of sintered density for bimodal powder mixtures. Metall Trans A 23(5):1455–1465. https://doi.org/10.1007/bf02647329 4. Geldart D, Abdullah EC, Verlinden A (2009) Characterisation of dry powders. Powder Technol 190(1–2):70–74. https://doi.org/10.1016/j.powtec.2008.04.089 5. Santomaso A, Lazzaro P, Canu P (2003) Powder flow ability and density ratios: the impact of granules packing. Chem Eng Sci 58(13):2857–2874

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6. Furnas CC (1931) Grading aggregates-I. Mathematical relations for beds of broken solids of maximum density. Ind Eng Chem 23(9):1052–1058 7. Goltermann P, Johansen V, Palbol L (1997) Packing of aggregates: an alternative tool to determine the optimal aggregate mix. ACI Mater J 94(5):435–443

Part VI

Additive Manufacturing of Metals: Fatigue and Fracture III

About a Digital Twin for the Fatigue Approach of Additively Manufactured Components Rainer Wagener, Matilde Scurria and Thilo Bein

Abstract A digital twin is an image of the reality. In case of a digital twin for additively manufactured components by the use of selective laser melting and Inconel® 718 issues as microstructure, anisotropy and load history are discussed. In order to take the main influences on the cyclic material behavior into account, local stresses and strains of a structure element are introduced. Furthermore, this new interpretation of the local load properties enables a two-stage combined macroscopic and microscopic material image as the basis for a digital twin for the fatigue assessment of additively manufactured components. Keywords Digital twin · Cyclic stress–strain behavior · Structure element

Nomenclature b c E I K’ n’ Ni εf σf X, Y, Z

fatigue strength exponent fatigue ductility exponent Young’s modulus Index of the fatigue regime, section of the tri-linear strain–life curve Cyclic hardening coefficient Cyclic hardening exponent Number of cycles to crack initiation Fatigue ductility coefficient Fatigue strength coefficient Directions linked to the process direction

R. Wagener (B) · M. Scurria · T. Bein Fraunhofer Institute for Structural Durability and System Reliability LBF, Bartningstr. 47, 64289 Darmstadt, Germany e-mail: [email protected] © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_34

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Introduction The development of additively manufactured structures and products is performed virtual from the first sketch to the printed component. In order to get a high-quality product that can be used for safety components under cyclic loading, the fatigue assessment including the structure optimization with respect to the exploitation of the lightweight design potentials must be done numerically. Therefore, the local component-related material behavior has to be available and considered. During the additive manufacturing process, multiple influences on the cyclic material behavior are caused, for instance, by the exposure strategy and surface conditions lead with the orientation of the loading to the building direction to mutual interactions. On the other hand industrial needs of low development time and numerical effort, the requirement on a digital twin for the fatigue assessment is growing and only accomplishable with an adjusted numerical concept with optimized material parameters.

Digital Twin for Fatigue Approach A digital twin refers to a digital replica of physical assets, processes and systems that can be used for various purposes [1]. The digital representation provides both the elements and the dynamics of how a component or a system behaves throughout its life cycle [2, 3]. In terms of a fatigue approach and besides the geometry of the component, the main influences on the cyclic material behavior as well as the loading conditions including the load time history must be represented by a digital twin. Furthermore, additively manufactured structures are well known for the heterogeneous microstructure depending on the exposure strategy and support structures. Due to these inhomogeneous microstructures, a local fatigue concept like the local strain concept or the material based fatigue approach [4] should be the first choice for the fatigue design of additively manufactured products. Hence, the foundation of a digital twin should be a local strain base fatigue approach concept, which enables the description of the cyclic material behavior by the use of a cyclic stress–strain and a strain–life curve. The choice of the strain–life curve depends on the scope of the fatigue approach. In case of a fatigue approach focused on the low cycle fatigue regime the strain–life curve according to Coffin–Manson–Basquin–Morrow [5–8] is commonly used Eq. 1. 

εa,t  εa,e + εa, p 

σf E



· (2Ni )b + ε f · (2Ni )c

(1)

With respect to growing computational power, the strain-based concept getting more and more into the focus even for the fatigue approach up to the very high cycle fatigue. Therefore, the fatigue life curve [9, 10], Eq. 2 has been developed in order to derive a continuous Wöhler curve from the Low Cycle Fatigue up to the very high

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cycle Fatigue regime. The formulation of a sum of the elastic and plastic strain part has been taken over from the Coffin–Manson–Basquin–Morrow rule, in doing so the elastic strain–life curve has been divided into three parts according to the three fatigue regimes of low cycle Fatigue, high cycle Fatigue and very high cycle Fatigue with the index I equal 1, 2 and 3. 

εa,t  εa,e + εa, p 

σ f,I E



· (2Ni )b I + ε f · (2Ni )c

(2)

Due to the setup of both strain–life curves, the derivation of the cyclic stress–strain curve according to Ramberg–Osgood [11], Eq. 3, are enabled. σa  σa  n1 + E K

εa,t  εa,e + εa, p 

(3)

The link between both, the stress–strain and the strain–life curve, is given by the so called compatibility conditions Eqs. 4 and 5. In case of the Fatigue Life curve, the material properties of the first regime have to be used. n 

b c

(4) 

K 

σf  

ε fn

(5)

In addition to the strain-controlled fatigue tests with constant amplitudes, which have to be performed to derive a strain–life curve, the incremental step test, which has been introduced by Landgraf [12] in 1968, can be used to derive the stress–strain curve with only one specimen. For that purpose, the load sequence of the incremental step test consists of subsequences of increasing and decreasing amplitudes. The resulting stress–strain behavior correlates better with the stress–strain behavior of randomly loaded structures than the stress–strain behavior which has been derived from the cyclic test with constant strain amplitudes [13, 14]. And last but not least, the strain-based fatigue approaches require only one strain–life curve, because different mean stresses and strains are considered by the use of damage parameters [15–21]. In comparison to stress-based fatigue approaches, the amount of fatigue test is reduced at the expense of numeric afford, because there is no need for Wöhler curves of different mean stresses or of different notch factors. Considering the requirement of the fatigue approach concepts, a digital twin should base on a strain-based concept. The basis for a reliable numeric fatigue approach is the description of the local stress–strain behavior [22]. Due to this reason, the main influences on the cyclic stress–strain behavior have to be identified and considered in case of a digital twin. Additively manufactured structures are well known for their heterogeneous microstructure with anisotropic behavior depending on the exposure strategy,

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Fig. 1 Coordinates

residual stresses and a rough surface. Furthermore, defects like pores and cracks could occur. Finally, the impact of the load time history on the cyclic material behavior has to be considered.

Test Setup For, the discussion of the different influences on the material behavior depending on the orientation to the additive manufacturing process direction the coordinate system according to Fig. 1 is used. The direction of powder deposition is X and the building direction Z. In case of additive manufacturing process of selective laser melting SLM with respect to process velocity, a subsize geometry for a flat specimen has been developed, Fig. 2. The transition zone from the highly load area to the clamping area has been optimized in order to decrease the notch factor. Furthermore, it is possible to build the specimens in build direction without support structures. To carry out strain-controlled fatigue tests with subsized specimens, the e-cylinder test rigs were developed, Fig. 3. Due to the mechanical setup cyclic tests with constant and variable amplitudes with low frequency enabled with a high accuracy and repeatability. In order to prevent buckling under compression loading an antibuckling device is used. For the purpose of analyzing the influences on the material behavior a print job with different specimen orientations and surface conditions has been designed. In order to analyze the cyclic material behavior, the Incremental Step Test is used. In doing so the orientation of the loading direction to building direction, different surface conditions and load capacity has been varied.

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Fig. 2 Subsize specimen geometry for fatigue tests

Fig. 3 E-cylinder test rig for strain-controlled fatigue test and a closed-up of the specimen with fixing of the extensometer and anti-buckling device

Stress–Strain System The basic idea of local fatigue concepts is to describe the local homogeneous material behavior by means of an infinitesimal material volume. In case of significant different microstructures and property gradients another set of material properties is required. With additively manufactured structures in mind, it is very difficult to define an infinitesimal material volume as well as specimens for the experimental derivation of the material properties even by the use of subsized specimen geometries, because additively manufactured structures contain a lot of different microstructures. With respect to numerical fatigue approach the effort will increase very strong. Having a closer look on additively manufactured structures, Fig. 4, different microstructures are visible. Due to the exposure strategy, the microstructure of the rim seems smoother and more homogeneous than the bulk material. The transition

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Fig. 4 Different microstructures of an Inconel® 718 component

could be interpreted as a metallurgical notch effect, which leads to a subsurface stress concentration and reduces the fatigue life. Even if the support structures are removed, the component behavior could still be influenced due to the local mass concentration and resulting cooling conditions as well as the rough surface. Depending on and caused by the microstructure, a digital twin requires the opportunity to deal with the heterogeneous material properties and rough surfaces. Due to the surface conditions, it could be difficult to measure a cross section in order to define a nominal stress even in the case of unnotched test specimens. On the other hand with reproducible surfaces as a function of the exposure strategy and support structures, it is sufficient to define a structure element with a technical surface. The resulting material or better structures properties contain the influence of these specified surface conditions. The classical way to consider irregularities like pores, surfaces or different microstructures caused by the manufacturing process during a fatigue approach is to start with a homogeneous material followed by influence factors to adjust the fatigue strength or life. Normally, the scope of a fatigue approach is not the local material behavior but the structure behavior. Keeping this in mind a transition from local stresses and strains to stresses and strains of a structure element should be useful

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for a digital twin for a fatigue approach of an additively manufactured structure. According to this the stresses and strains, as well as the Young’s modulus, have to be interpreted as a resulting stress, strain or stiffness of structure elements. Well knowing that due to stress concentrations and irregularities the local stresses and strains could be higher.

Cyclic Stress–Strain Behavior After removing the specimens from the building plate, an ultrasonic cleaning followed by a stress relieving heat treatment. Therefore the specimens were heated up to 650 °C in about 9 h with to stops at 250 and 500 °C. After a isotherm last for 2 h, the heat treatment was continued by cooling down to 250 °C in 4 h and finally, rapid to room temperature under gaseous atmosphere. A part of the specimens was left with this conditions and the other part was subjected to abrasive blast cleaning. For all the different material conditions, Incremental Step Tests with the maximum load of εa,t  0.8%, εa,t  0.6% and εa,t  0.4% were carried out. Depending on the maximum load, the cyclic stress–strain behavior differs between elastic–plastic stress-strain behavior in case of εa,t  0.8% and more or less macroscopic linear elastic stress–strain behavior in case of a maximum load of εa,t  0.4%, Fig. 5. The load maximum of εa,t  0.6% is comparable to the εa,t  0.8%. The amount of cyclic softening depends on the load capacity. Below the focus of the discussion will be on the cyclic material behavior with a maximum load of εa,t  0.4%, because in this case only a small amount of plasticity occurs at all. Therefore, the stabilized cyclic stress–strain curves are derived using the measured stresses and strains at the middle of the fatigue life. The resulting charts are printed in Fig. 6 and summarized in Table 1. The cyclic stress–strain curve described by the reversal points does not hit the origin of the coordinates necessarily. Especially, the stress–strain curves of specimens which were built perpendicular to the cyclic loading direction have an offset in stress direction. One of the reasons for the behavior could be the influence of the removed support structure which is required to build the specimens and is positioned along the lower side. Besides, the irregular geometry the microstructure could be influenced by the different cooling curves due to the former mass concentrations. Up to now the influence of the support structures on the resulting material behavior is still under research. Mean stresses are regarded to influence the fatigue life. Tensile stresses lead to a reduced fatigue life while compression stresses lead to an increased fatigue life. This well-known relationship is also valid for additive manufactured structures such as visible from the comparison of the mean stresses, Fig. 7, and the fatigue life, Fig. 8. Regarding the X direction with increasing mean stresses, the fatigue life is reduced and could be influenced by the surface condition and additional heat treatment. With respect to a digital twin for the fatigue approach of additively manufactured

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Fig. 5 Development of the cyclic stress–strain behavior of selective laser melted Inconel® 718 with respect to the load capacity

components, it is not satisfying to describe only one cyclic stress–strain curve. In fact, the digital twin has to deal with the different orientations as well as load history.

Conclusions and Outlook The microstructures of additively manufactured structures are depending on the exposure strategy. Due to the resulting inhomogeneous microstructures with property gradients, a local strain-based concept for the fatigue approach should be included in a digital twin for the numerical fatigue approach. The macroscopic view using stresses and strains of a structure element instead of an infinitesimal material volume reduces the numerical effort and enables two ways to derive cyclic structure properties. Performing cyclic tests with additively manufactured specimens like Fig. 2 enable the derivation of macroscopic structure properties. The second way is to start with the microscopic scale and build up the structure element if necessary. Doing so, the numerical effort could be divided into two parts and performed separately. The digital twin for a fatigue approach of additively manufactured and cyclic loaded struc-

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Fig. 6 Cyclic stress–strain curves for a maximum load capacity of εa,t  0.4% at half number of cyclic to crack initiation Fig. 7 Occurring mean stress in incremental step tests with a total strain amplitude of εa,t  0.4%

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Table 1 Results of incremental step test with a maximum load of εa,t  0.4% Direction

Upper side

Lower side

Blasted/not Mean blasted stress σm (MPa)

Young’s modulus E (GPa)

X

As build

Polished

Not blasted

380

188

16320

Polished

Support structure

Not blasted

−12

147

40800

Polished

Support structure

Blasted

435

155

4720

Polished

Polished

Not blasted

1

184

115040

As build

As build

Blasted

34

168

46720

Polished

Polished

Not blasted

−9

168

45200

As build

Support structure

Not blasted

−93

198

20160

As build

Support structure

Blasted

−56

202

24400

Polished

Polished

Not blasted

54

219

33600

Z

XZ

Fig. 8 Fatigue Life for a total strain amplitude of εa,t  0.4%

Number of cycles to crack initiationNi

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tures consists of a two-stage process. The first stage deals with the microstructure (microscopic point of view) and allows calculating the behavior of the macroscopic structure elements, which could be derived by the cyclic tests, too. Acknowledgements The research and development project ‘BadgeB’ that form the basis for this publication is funded within the scope of the “Additive Fertigung—Individualisierte Produkte, komplexe Massenprodukte, innovative Materialien” by the Federal Ministry of education and Research and managed by the KIT project management agency “Projekträger Karlsruhe—Produktion und Fertigungstechnologien”. The authors are responsible for the content of this publication.

References 1. Minds + Machines: Meet A Digital Twin, Youtube, GE Digital, retrieved 26 July 2017 2. Digital Twin (2018) Wikipedia.org. Accessed 19 Aug 2018 3. Introduction to Digital Twin (2017) Simple, but detailed, Youtube. IBM Watson Internet of Things. Accessed 27 June 2017 4. Hell M, Wagener R, Kaufmann H, Melz T (2017) Strategies for material modelling regarding fatigue design under variable amplitude loading with strain-based fatigue design approaches. In: Proceedings of the 7th international conference on fatigue design, 2017, CETIM. ISBN 978-2-36894-126-3 5. Coffin Jr LF (1954) A study on the effect of cyclic thermal stresses on a ductile metal, Trans. ASME 76:931–950 6. Manson SS (1965) Fatigue: a complex subject—some simple approximations. Exp. Mech. 5(7):45–87. https://doi.org/10.1007/bf02321056 7. Morrow JD (1965) Cyclic plastic strain energy and fatigue of metals. ASTM STP 278:45–87. https://doi.org/10.1520/STP43764S 8. Basquin OH (1910) The exponential law of endurance tests. Proc ASTM 10:625–630 9. Wagener R, Melz T (2017) Deriving a continuous fatigue life curve from LCF to VHCF, SAE 2017-01-0330. https://doi.org/10.4271/2017-01-0330 10. Wagener R, Melz T (2018) Fatigue life curve—a continuous wöhler-curve from LCF up to the VHCF regime. Mater. Test. 60(2018):10. https://doi.org/10.3139/120.111233 11. Ramberg W, Osgood WR (1943) Description of stress-strain curves by three parameters. National advisory committee for aeronautics technical note no 902 12. Landgraf WR, Morrow JD, Endo T (1969) Determination of the cyclic stress-strain curve. J Mater 4:176–188 13. Christ H-J (1998) Materialermüdung und Werkstoffmikrostruktur, Hrsg H-J Christ, WerkstoffInformationsgesellschaft, Frankfurt 14. Wagener R (2007) Zyklisches Werkstoffverhalten bei konstanter und variabler Beanspruchungsamplitude, PhD-thesis, TU Clausthal 15. Smith KN, Watson P, Topper TH (1970) A stress-strain function for the fatigue of metals. J Mater 5(4), S 767–778 16. Bergmann J (1983) Zur Betriebsfestigkeitsmessung gekerbter Bauteile auf Grundlage der örtlichen Beanspruchungen. Dissertation, Technische Universität Darmstadt (1983) 17. Haibach E, Lehrke HP (1975) Das Verfahren der Amplituden-Transformation. Fraunhofer LBF, Darmstadt, Bericht Nr FB-125 18. Heitmann HH (1981) Vorhersage der technischen Anrißlebensdauer unter Berücksichtigung des Verhaltens von Mikrorissen. Dissertation, TH Aachen 19. Werner S (1999) Zur betriebsfesten Auslegung von Bauteilen aus AlMgSi 1 unter Berücksichtigung von hohen Mitteldehnungen und Spannungskonzentrationen. Dissertation, Technische Universität Darmstadt, Fachbericht FB-217, Darmstadt

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20. Best R (1991) Der Schadensparameter im Kerbgrundkonzept. Materialprüfung 33(6):184–188 21. Vormwald M (1989) Anrißlebensdauervorhersage auf Basis der Schwingbruchmechanik für kurze Risse. Dissertation, Technische Hochschule Darmstadt, Darmstadt 22. Hell M, Wagener R, Kaufmann H, Melz T (2017) Effect of the material modeling and the experimental material characterisation on fatigue life estimation within strain-based fatigue assessment approaches. In: FATIGUE 2017—the 7th engineering integrity society international conference on durability and fatigue, EIS. Cambridge

Effect of the Surface Finish on the Cyclic Behavior of Additively Manufactured AlSi10Mg Matilde Scurria, Benjamin Möller, Rainer Wagener and Tobias Melz

Abstract The design flexibility offered by the newest additive manufacturing technologies is attracting the attention of the automotive industry for the realization of safety-relevant components. However, the realization of complex geometries is characterized by the use of support structures which sustain surfaces with downskin angles lower than 45°. The subsequent removal of these punctual joints leaves the surface irregular and with a large amount of defects. In this work, the effect of surface imperfections on the cyclic stress–strain behavior of additively manufactured metals is evaluated. Small-scale specimens are manufactured by selective laser melting of AlSi10Mg powder. The specimens are manufactured using different build orientations, part of them are left as-built while the surface of other specimens has been polished. Incremental step tests are carried out in order to evaluate the cyclic stress–strain behavior of this material. Keywords SLM · AlSi10Mg · Additive manufacturing · Cyclic material behavior · IST · Fatigue

Introduction The idea of a new technology that enables the production of extremely complex geometries is pushing the research above the limit represented by rapid prototyping towards the additive manufacturing of metals. The new challenge is the use of these technologies for the production of safety-relevant metal components. While the reproducibility and stability of the processes develop fast, due to the research performed by the machine manufacturers, a method for a fatigue assessment of addiM. Scurria (B) · T. Melz Research Group of System Reliability, Adaptive Structures and Machine Acoustics SAM, Technische Universität Darmstadt, Darmstadt, Germany e-mail: [email protected] B. Möller · R. Wagener · T. Melz Fraunhofer Institute for Structural Durability and System Reliability LBF, Darmstadt, Germany © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_35

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tively manufactured components is still missing. In fact, the variables which can affect the cyclic material behavior are several. Most of the times, they are reciprocally connected, and therefore hard to isolate. Due to this fact, a material-based fatigue approach [1] is to be preferred. In this method, the S-N curves of the component are derived using a local strain approach instead of a nominal-stress based one, where the effects of the different variables are evaluated by carrying out a larger number of load-controlled tests. In case of a local strain approach, the experimental campaign is limited to the evaluation of the cyclic stress–strain behavior, derived by just one strain-controlled cyclic test with variable amplitude, and to one strain–life curve. In this work, the effects of anisotropy, heat treatment and surface finish on the cyclic stress–strain behavior on AlSi10Mg [2, 3], an aluminum alloy close to the eutectic composition and commonly used by the casting industry, produced by Selective Laser Melting (SLM) [4], are investigated.

Experimental Campaign Preparation of the Specimens The nominal geometry of the subsize specimens for fatigue testing is represented in Fig. 1 (left). To describe the effect of the anisotropy, specimens oriented along the powder deposition direction in X (0°, lying), XZ (45°) and Z (90°, standing) directions and specimens lying orthogonal to the powder deposition direction with respect to the building platform (Y specimens) are manufactured using standard process parameters (Fig. 1, right). Figure 1 (center) additionally shows slides of SLM material for manufactured and subsequently polished specimens of the same geometry. For both as-built and polished specimens, half of them are subjected to stress relief heat treatment. In this way, each of the four build directions X, XZ, Z, and Y has four states according to the table of Fig. 1: as-built, tempered and not polished, not tempered but polished and tempered and polished. For each one of these sixteen configurations, three specimens are produced, being tested with three different maximal strain amplitudes of εa,t  0.4, 0.6 and 0.8%. Individually adapted support structures, needed for lying (X and Y specimens) and 45° XZ specimens or rather surfaces with downskin angles up to 50° [5], are subsequently mechanically removed. The effect of support structures has to be taken into account, since, even after their removal, the material in the presence of these punctual joints and of its substrate shows several defects. In Fig. 2a, the embedded section cut of half of a specimen built in the X-direction is represented. Figure 2b shows the part of material which is in contact with the support structures, where several pores can be found. As shown in Fig. 2b, even larger pores have been detected in the area of the melted contour covered by the laser before starting to fill the core material (exposure strategy) (Fig. 2c). The irregularity of the material is already obvious by the analysis of microsections using a light microscope at low magnification. After

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385 Build direction X

XZ Z

yes no yes no yes no yes no

Tempered

Y

Polished yes no 3 3 3 3 3 3 3 3 3 3 3 3 3 3 3 3

Fig. 1 Geometry of the specimens (left), different build directions within the build platform (center) and configurations for fatigue testing (right)

Fig. 2 Section cut of an embedded specimen built in the X-direction (a), analysis under the light microscope of the area close to the support structures (b) and of the contour (c)

the production, the specimens are removed from the build platform. The Z specimens are ready for fatigue testing, while the X, Y and XZ specimens are subjected to a subsequent machining for the removal of support structures. The surfaces reveal that they were previously in contact with the punctual joints and their roughness results in higher values of Rz  55–120 μm than the other AM surfaces due to the large amount of pores and defects. For AM surfaces without support structures, typical values of the roughness are Rz  19–24 μm. Another difference between the upskin and downskin surfaces in dependency on support structures is given by the cooling rate. The upskin surfaces, skimmed by the air, are subjected to a faster cooling rate compared to the downskin ones. Some studies have already shown how a correct design of the support structures can reduce differences in terms of residual stresses and microstructures [6–8]. However, this aspect is not the object of this study and will be addressed to further investigations. The specimens extracted from the slides are mechanically polished to satisfy a final roughness of Rz ≤ 2 μm.

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Fig. 3 E-Cylinder test rig used for the incremental step tests (left), test setup (center) and example of load sequence [9] (right)

Fig. 4 Failure envelope (left), reversal point of one block (center) and their use for the evaluation of the cyclic stress–strain curve through the Ramberg–Osgood equation (right)

Test Setup The experimental campaign is carried out using an E-cylinder test rig, developed at the Fraunhofer LBF, which enables strain-controlled fatigue testing using an electric motor. A load cell measures the force, while the strain is controlled by an extensometer. An anti-buckling device is used to avoid buckling when a strain ratio of Rε < 0 is applied. The test rig and the setup are shown in Fig. 3 (left). The specimens are subjected to Incremental Step Tests (IST) [10, 11], where a defined load sequence with decreasing and increasing strain amplitude, at a strain ratio Rε  −1, is imposed, until the failure occurs. The cycles between two maxima belong to one block (Fig. 3, right). The failure corresponds to the n-block, in which the maximum force decreased by 10% relates to n/2, which is called the stabilized block or stabilized state (Fig. 4, left). The reversal points (Fig. 4, center) of the stabilized block are used to evaluate the cyclic stress–strain curve by regression of the Ramberg–Osgood equation [12] (Fig. 4, right).

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Table 1 Results obtained for ISTs with a maximal strain amplitude of εa,t  0.4% Build direction

Polished

Tempered

E [GPa]

X

No

No

Yes No Y

XZ

Z

Mean stress σm [MPa]

Cycles to crack initiation Ni

72

10

1280

No

85

50

6960

Yes

72

20

1600

Yes

Yes

70

17

5280

No

No

80

25

640

Yes

No

80

30

3520

No

Yes

70

10

1393

Yes

Yes

72

−10

16,880

No

No

80

5

3440

Yes

No

80

80

2720

No

Yes

70

0

3680

Yes

Yes

70

30

5600

No

No

80

0

7600

Yes

No

75

32

14,880

No

Yes

70

0

3760

Yes

Yes

70

15

11,120

Test Results The incremental step tests are performed at three different maximal strain amplitudes εa,t  0.4, 0.6 and 0.8%. However, only the results of the tests performed with εa,t  0.4%, summarized in Table 1, will be discussed in this study, since the plastic components are still limited and the cyclic stress–strain behavior for the different build directions differs the most. The intersection between the curves and the stress axis may assume values different from zero, and this value is indicated in Table 1 as mean stress σm .

Effects of a Stress Relief Heat Treatment In Fig. 5, the stabilized cyclic stress–strain curves for both tension and compression of each of the four build directions X, Y, XZ and Z are represented. The curves are evaluated separately for the first (tension) and third (compression) quadrant, since they may slightly differ. For all four built directions, the specimens in the polished

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Fig. 5 Stabilized stress–strain curves in tension and compression for a maximal strain amplitude of εa,t  0.4% for the four built directions X (a), Y (b), XZ (c) and Z (d)

and not tempered state show a linear–elastic behavior, while for the configuration XZ and Z also the as-built state is linear–elastic. In general, the tempered state has a larger plastic component. The cyclic material behavior is influenced by the heat treatment, while, as expected, is not affected by the surface finish. Especially in case of tempered Z specimens, the more regular surface, but not polished, does not affect the cyclic stress–strain curve. If the surface finish is not influencing, the cyclic behavior, on the other hand, is affecting the number of cycles to crack initiation, whose evaluation is not the aim of the ISTs, the IST can be used, for example, to understand the mechanisms that cause the fatigue failure.

Anisotropy In order to analyze the effect of the anisotropy, for each of the four states of treatment (as-built, not tempered but polished, tempered but not polished and tempered and

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polished) the results of the tests performed on specimens built along different orientations are compared. The highest anisotropic behavior belongs to the specimens in the as-built state (Fig. 6a). In this case, together with a variation of the Young modulus, another important difference can be observed: the plastic component of the total strain increases significantly for specimens oriented from X to Z and XZ, respectively. This means that the as-built state may be critical for the simulation phase, where the material cannot just be described by the four parameters of Young modulus, Poisson ratio, cyclic hardening coefficient K’ and cyclic hardening exponent n’ due to their dependence to the build direction. Furthermore, even in this case, the cyclic material behavior of the X and Y specimens is represented by the same curve, making these two directions indistinguishable and reducing the number of variables represented by the build orientation from the original four to three (X, XZ and Z). For the not tempered, but polished state, the behavior is linear–elastic with a slightly varying Young modulus (Fig. 6b). For the tempered case, both the not polished (Fig. 6c) and polished (6d) state, the same courses of the curves appear. This means, if the material undergoes a stress relief heat treatment, the number of variables decreases further to the simple situation where neither the influences of build orientation nor of the surface finish occur, so that the cyclic material behavior is adequately described by one single stress–strain curve. In addition, the course in tension and compression is completely symmetric and therefore the cyclic stress–strain curve can be evaluated based on amplitudes using all the reversal points mirrored in the first quadrant, as illustrated in Fig. 4 (right), without losing important information.

Cycles to Crack Initiation As already mentioned, the surface finish has a minor influence on the cyclic material behavior, but it is influencing the fatigue life. In Fig. 7, the number of cycles to crack initiation of all sixteen configurations is represented. The effect of the surface preparation is evident for all the cases. With a single exception for the not tempered (NT), polished state of the XZ specimen, ISTs on polished specimens show a larger fatigue life compared to the as-built (AB) state. The fracture mechanisms will be discussed in the paragraph regarding the metallographic investigations and could explain this found anomaly exception. Especially for the specimens built in the X- and Y-directions, the beneficial effect of the surface polishing is evident and is related to the statistical nature of fatigue failures, which occur in proximity of defects or inclusions.

Metallographic Investigations Some part of the AM material and the fracture surfaces have been analyzed to understand the internal structure of the material and the phenomena that caused the failure

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Fig. 6 Stabilized stress–strain curves in tension and compression for a maximal strain amplitude of εa,t  0.4% for the four states: as-built (a), not tempered but polished (b), tempered but not polished (c) and tempered and polished (d) Fig. 7 Cycles to crack initiation for different combinations of the build direction, heat treatment and surface finish

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391

Fig. 8 Microsections of the areas adjacent to the specimens manufacture

of the specimens. First, the area of the slides in correspondence of the specimens has been analyzed, to state the initial material conditions (Fig. 8). It can be noticed from the microsections of Fig. 8 that the considerable thickness of the contour can reach the value of 0.5 mm in some points and therefore requires high precision and time to remove the entire contour layer just by polishing. A total removal of about one millimeter would be necessary to avoid effects of the boundary area, where the microstructure is different and the large amount of defects can easily cause the initiation of cracks, on the fatigue fracture. In addition, it has to be underlined that polishing is not only unsuitable to reach the core of the material, but also that the removal of a thin layer of material reveals imperfections present in the substrate, e.g. pores, which are geometrical notches for the stress field of the material, leading to stress concentrations. Another important consideration to be taken into account is that the specimens can show differences from the polished sides to the polished surfaces. While the sides consist of core material without the characteristics of the contour layer, the critical areas may be found on the polished surfaces of the specimens, where just a thin layer of the contour is removed, revealing the defects of the substrate. In fact, it was observed that the crack starts if not polished samples are always at the sides of the specimen or at its edges, while for polished specimens the failure occurs also due to defect located on the larger surfaces in the center of the specimens, where the contour layer is still present.

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Some examples of fracture surfaces analyzed with the use of an SEM microscope are shown in Fig. 9. The SEM images of Fig. 9a, c illustrate the fracture surface of two polished specimens. It can easily be distinguished between the core area and the contour layer, which shows imperfections. In case of the specimen of Fig. 9a, the crack starts from the edge with an abundance of defects (Fig. 9b). On the other hand, Fig. 9c shows the case in which the crack starts from a single pore present on the larger surface of the specimen. Figure 9e shows the fracture surface of an as-built specimen, where the two areas of the core material and of the contour layer can easily be distinguished. The contour layer also includes the sides of the specimen.

Conclusions In this work, the cyclic stress–strain behavior of additively manufactured AlSi10Mg in case of as-built, tempered but not polished, polished but not tempered and both polished and tempered states with respect to different built directions has been investigated. The stress relief heat treatment causes an increase in the plastic strain component in the cyclic stress–strain curve compared to the not tempered, e.g. as-built, state, already for a maximal total strain amplitude of εa,t  0.4%. The heat treatment is affecting the material behavior, while the surface finish has just a slight influence, and only for X and Y specimens, where the surfaces are highly irregular before the polishing process. An anisotropic material behavior from linear–elastic for the asbuilt state to elastic–plastic for polished states is shown within the different build directions and will be challenging when deriving an appropriate material model. But it becomes isotropic as soon as the material is subjected to a stress relief heat treatment. These considerations are very helpful when the material is characterized during the first stage of material-based fatigue approach. Furthermore, as usual, the process of mechanical polishing increases the number of cycles to crack initiation of the material. This effect might be traced back to both the introduction of compressive residual stresses on the surface of the material and to the positive effect of the significantly reduced roughness of the surface. However, the second aspect can be improved by the removal of the entire contour layer to reveal the core material on the specimens’ surface.

Outlook The further investigations aim a better understanding of the influences of support structures and of the contour area of the material produced by additive manufacturing. On the other hand, main attention will be turned on the fatigue life estimation using local strain approaches and an implementation of the achieved results for the cyclic material behavior into the numerical simulation.

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Fig. 9 Overview of the fracture surfaces for polished specimens (a, c) and details showing the primary crack initiation (b, d) for polished specimens, and for an as-built specimen (e)

Ackowledgements The research and development project “VariKa” that forms the basis for this report is funded within the scope of the “PAiCE Digitale Technologien für die Wirtschaft” technology programme run by the Federal Ministry for Economic Affairs and Energy and is managed by the DLR project management agency “Gesellschaft, Innovation, TechnologieInformationstechnologien/Elektromobilität” at the German Aerospace Center in Cologne. The author is responsible for the content of this publication.

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References 1. Hell M et al (2015) Fatigue life design of components under variable amplitude loading with respect to cyclic material behaviour. Proc Eng 101:194–202 2. Brandl E et al (2012) Additive manufactured AlSi10Mg samples using selective laser melting (SLM): microstructure, high cycle fatigue, and fracture behavior. Mater Des 34:159–169 3. Kempen K et al (2012) Mechanical properties of AlSi10Mg produced by selective laser melting. Phys Proc 39:439–446 4. Wilhelm Me (1999) Direktes Selektives Laser Sintern einkomponentiger metallischer Werkstoffe, Dissertation, RWTH Aachen 5. V D I. 3405: Blatt 3 (2015) Additive Fertigungsverfahren – Konstruktionsempfehlungen für die Bauteilfertigung mit Laser-Sintern und Laser-Strahlschmelzen. VDI-Gesellschaft Produktion und Logistik 6. Mishurova, T et al (2018) The influence of the support structure on residual stress and distortion in SLM Inconel 718 parts. Metall Mater Trans A 49.7:3038–3046 7. Calignano F (2014) Design optimization of supports for overhanging structures in aluminum and titanium alloys by selective laser melting. Mater Des 64:203–213 8. Poyraz Ö et al (2015) Investigation of support structures for direct metal laser sintering (DMLS) of IN625 parts. In: Proceedings of the solid freeform fabrication symposium, Austin, Texas, USA (2015) 9. Masendorf R, Wagener R. Prüf- und Dokumentationsrichtlinie für die experimentelle Ermittlung mechanischer Kennwerte von Feinblechen aus Stahl für die CAE-Berechnung, Appendix Incremental Step Test 10. Morrow J (1965) Cyclic plastic strain energy and fatigue of metals. Internal friction, damping, and cyclic plasticity. ASTM International (1965) 11. Landgraf RW, Morrow J (1969) Determination of the cyclic stress-strain curve. J Mater 4(1):176–188 12. Ramberg W, Osgood WR (1943) NACA technical note No. 902

Effect of Heat Treatments on Fatigue Properties of Ti–6Al–4V and 316L Produced by Laser Powder Bed Fusion in As-Built Surface Condition Antonio Cutolo, Chola Elangeswaran, Charlotte de Formanoir, Gokula Krishna Muralidharan and Brecht Van Hooreweder

Abstract Over the last decade, additive manufacturing (AM) techniques have been expanding rapidly due to their ability to produce complex geometries with an efficient use of material. In order to design reliable AM parts, the mechanical properties resulting from the manufacturing process need to be understood. The present study investigates the fatigue of AM Ti–6Al–4V and 316L. Miniaturized Ti–6Al–4V and 316L specimens were manufactured using laser powder bed fusion (L-PBF). The geometry, process parameters, and loading conditions were kept constant and the specimens were tested in as-built surface condition. The S-N curves of as-built, stressrelieved and HIP’ed specimens were measured, and an analysis of the microstructure, relative density and surface roughness was performed. The effect of fatigue influencing factors (residual stresses, surface roughness, porosity and microstructure) was systematically investigated. In order to understand the fatigue failure mechanism, identification of crack initiation point, via fracture surfaces analysis, was performed. Keywords Laser powder bed fusion · Ti–6Al–4V · 316L · Fatigue · Heat treatment

Introduction Laser powder bed fusion (L-PBF), also known as selective laser melting (SLM), is an additive manufacturing (AM) technique that allows the production of metallic parts using a high-power laser that locally melts successive layers of powder. The L-PBF

A. Cutolo (B) · C. Elangeswaran · C. de Formanoir · B. Van Hooreweder Department of Mechanical Engineering, KU Leuven, Celestijnenlaan 300, 3001 Leuven, Belgium e-mail: [email protected] A. Cutolo · C. Elangeswaran SIM M3 Program, Technologiepark 935, 9052 Zwijnaarde, Belgium G. K. Muralidharan 3D Systems Leuven, Grauwmeer 14, 3001 Leuven, Belgium © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_36

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technique has several advantages over conventional production techniques, such as a low material waste, a high level of flexibility and the ability to produce geometrically complex structural parts. The details and applications of this AM process have been widely reviewed [1, 2]. Due to their ability to fabricate customized parts with complex geometry, LPBF processes are increasingly used in many different industrial sectors, such as aerospace, biomedical and automotive industry. However, one of the major drawbacks of powder bed AM processes such as L-PBF is the very high surface roughness induced by the process [3]. This is particularly critical considering the fact that a customized SLM part, with a complex geometry, often has some areas with rough as-built surfaces, since in many cases, not all surfaces are accessible for polishing or machining. For that reason, this work focusses on the behavior of metal AM samples with as-built surface condition. In order to exploit the benefits of AM techniques for safety-critical component production, the mechanical behavior of additively manufactured metals must be better understood. For this reason, fatigue of AM’ed parts is widely investigated in the literature. Leuders et al. [4] concluded that for Ti–6Al–4V manufactured by SLM, the main drivers for fatigue crack initiation and failure are internal defects, such as pores, acting like stress raisers. Kasperovich and Hausmann [5] found that hot isostatic pressing (HIP), by coarsening the microstructure and reducing porosity, leads to a significant improvement of ductility and fatigue strength of Ti–6Al–4V processed by SLM. Therefore, they concluded that HIP is necessary for cyclically loaded components. Similar conclusions are reported by Ramulu and Edwards [6] who also indicated the need to perform heat treatments, i.e. stress-relief or HIP, on SLM Ti–6Al–4V in order to reduce the residual stress levels generated during manufacturing. In addition, they concluded that the effect of as-built surface condition on fatigue properties could not be determined due to the presence of internal defects and residual stress. Cao et al. [7] reported that fatigue properties of SLM Ti–6Al–4V are mostly controlled by surface roughness and defects that significantly reduce the fatigue strength of the material. Austenitic 316L exhibits very high ductility [8, 9] that helps reducing the local stress at defect sites and delaying crack initiation, hence reducing the influence of defects. This could explain the relatively good fatigue behavior of 316L processed by L-PBF [8]. Spierings et al. [10] tested the fatigue behavior of SLM 316L, using post-machined specimens, and observed that the material was not very sensitive to pores and other defects. From a microstructural point of view, grain coarsening and homogenization through recrystallization induced by dedicated heat treatments were found to improve crack initiation behavior [11]. Examination of the state of the art shows that fatigue properties of L-PBF Ti–6Al–4V and 316L are mostly influenced by surface roughness, residual stresses produced during manufacturing, microstructure and pores. In order to exclude the influence of porosity from the scope of this investigation, the fatigue performances of L-PBF Ti–6Al–4V and 316L were measured on almost pore-free specimens. A systematic investigation of the influence of heat treatments on fatigue properties was performed on these specimens.

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Fig. 1 Fatigue sample geometry and dimensions

Materials and Methods Fatigue specimens used in this investigation are miniaturized samples with a circular cross section and continuous radius between ends. These miniaturized test coupons were designed specifically for the fatigue testing of additively manufactured metals, taking into account the following considerations: (i) the radius between the two ends is optimized in order to reduce the impact of staircase effect on surface roughness [12]; (ii) the design ensures that failure always occurs in the smallest cross-sectional area, minimizing the scatter in fatigue data; (iii) buckling during axial compression loading is avoided; (iv) the gripping area is optimized for sufficient friction between the sample and the clamps; (v) the miniaturized geometry design reduces the production cost and time. Sample geometry and dimensions are shown in Fig. 1. Miniaturized samples were manufactured via L-PBF on a ProX DMP320 machine (3D Systems) equipped with a 500 W Fiber laser using LaserForm Ti Grade 23 powder (3D Systems) for Ti–6Al–4V specimens and Laser Form 316L powder (3D Systems) for 316L specimens. A layer thickness of 60 μm was used for both Ti–6Al–4V and 316L specimens. The test coupons were produced with the axis oriented in the building direction (i.e. Z axis). Two different heat treatments were investigated for Ti–6Al–4V samples, namely, stress-relief (SR) and hot isostatic pressing (HIP): for SR, the samples were heated at 850 °C for 2 h followed by air cooling, while for HIP the samples were heated at a temperature of 920 °C at a pressure of 1000 bar for 2 h. For 316L samples, only SR was performed, by heating the 316L samples to 470 °C for 5 h. In total, three different conditions were investigated for Ti–6Al–4V, i.e. as-built (ASB), SR, HIP and two conditions for 316L, i.e. ASB and SR. All the samples were tested in as-built surface condition. For each of the above-mentioned conditions, Archimedes’ density measurements were performed on all the samples. An Acculab AST 224 Sartorius balance with an accuracy of 0.1 mg equipped with a calibrated Sartorius Archimedes’ kit was used for gravimetric measurements. A Keyence VHX-6000 optical microscope was used for surface roughness evaluation. Optical microscopy was performed with Keyence VHX-6000 on ground, polished and etched samples. A2% HF solution was used to etch Ti–6Al–4V samples to detect possible alpha case if present. 316L samples were electrochemically etched at 6 V DC with 10% oxalic acid solution. Tensile tests were performed on an SIMADZU AG-X Plus machine according to ASTM E8M standard for quasi-static properties evaluation of the two materials

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Table 1 Mechanical properties, density and surface roughness of Ti–6Al–4V and 316L samples for different post-treatments Ti–6Al-4V

316L

Heat treatment

ASB

σ y (MPa)

1086 ±15

926 ±45

860 ±40

453 ±7

440 ±20

σU T S (MPa)

1246 ±22

1002 ±30

960 ±30

573 ±6

570 ±30

11 ±1.0

Elongation (%) ρr el density (%)

99.61 ±0.05

SR

HIP

16 ±3.0 99.63 ±0.06

18 ±4.0 99.71 ±0.14

ASB

46 ±1.0 99.42 ±0.06

Ra (μm)

5.13 ±0.61

7.23 ±1.30

Rv (μm)

13.79 ±2.09

17.90 ±1.86

Rz (μm)

29.21 ±3.33

38.14 ±5.87

SR

49 ±5.0 99.42 ±0.05

when different post-treatments were considered. An Instron Electropuls E10000 with a 10 kN load cell was used for all fatigue tests. The tests were force-controlled and were conducted with a constant amplitude fully reversed sinusoidal load (stress ratio R  −1) with a frequency of 60 Hz. Fatigue tests were performed on the five different batches until failure of the sample or stopped after 2 × 106 cycles. For each batch, at least three different stress levels were investigated and, for each stress level, at least two samples were tested. Table 1 classifies the samples in terms of material, heat treatment, quasi-static mechanical properties, density and surface roughness. A Philips XL 30 FEG scanning electron microscope (SEM) was used for crack initiation points identification on fractured samples.

Results and Discussion Relative Density The relative densities were measured by Archimedes’ method and are reported in Table 1. Ti–6Al–4V samples exhibit a low fraction of residual porosity with a relative density larger than 99.50%, even in as-built condition. The HIP post-process slightly reduces the amount of residual porosity for Ti–6Al–4V. For 316L samples, the measured relative density was slightly below 99.50% both in as-built and stress-relieved conditions. The process parameters used for sample production were standard parameter sets for LaserForm Ti gr 23(A). and LaserForm 316L(A). As a result, the amount of residual porosity generated during the manufacturing process was systematically below 0.6%.

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Fig. 2 Microstructure for Ti–6Al–4V—ASB (a), Ti–6Al–4V—SR (b), Ti–6Al–4V—HIP (c), 316L—ASB (d) and 316L—SR (e); cross section perpendicular to the building direction

Microstructure Characterization Ti–6Al–4V exhibits an acicular martensitic microstructure in the as-built condition. Fine α martensite needles can be seen in Fig. 1a. Prior columnar β grains were also observed along the building direction. Rapid cooling rates in SLM lead to such a typical process-induced martensitic microstructure [13]. A fine α + β lamellar microstructure is obtained after SR, as shown in Fig. 2b. This microstructure results from the decomposition of the α martensite to α + β. Following HIP, the lamellar α + β microstructure exhibits an even coarser morphology, as shown in Fig. 2c. 316L in as-built condition exhibits a typical cellular microstructure with clear distinction of melt pools as shown in Fig. 2d. A columnar grain growth pattern can also be observed, with grains growing across melt pool boundaries. Grain growth direction within the melt pool tends to follow the temperature gradient within the pool, starting from the periphery and growing towards the core. Stress-relief heat treatment did not induce any major microstructural change. A similar cellular microstructure is observed, as shown in Fig. 2e.

Quasi-static Properties Tensile properties of the two materials including yield stress, ultimate tensile strength and elongation at break are reported in Table 1. From these results, different observations can be made for Ti–6Al–4V and 316L: (i) in as-built condition Ti–6Al–4V exhibits high strength and relatively low elongation (11%); (ii) SR and HIP posttreatments improve elongation at break of Ti–6Al–4V while reducing ultimate and

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(a)

(b)

400

350

350

300

300

σ

σ

400

250

250

200

200

150

150

100 104

105

106

100

316 316

104

105

106

Fig. 3 Influence of post-treatments on fatigue properties of Ti–6Al–4V (a) 316L (b) in as-built surface condition

yield strength; and (iii) 316L in stress-relieved condition exhibits similar yield and ultimate strength to those measured for the as-built material, with a slight (3%) rise of the elongation at break. It is well known [14] that in as-built SLM’ed Ti–6Al–4V, the martensitic microstructure is responsible for its relatively high strength and low ductility. With stress-relief and, to a larger extent, with HIP, the α microstructure coarsens into α + β lamellae, leading to a reduction in ultimate and yield stress and an increase in ductility. The coarser the grains (Fig. 2b, c), the more pronounced the abovementioned effect, which can be inferred from lowest strength and highest ductility of HIP’ed samples. With 316L, the SR heat treatment performed in order to relieve SLM-induced residual stresses did not cause much microstructural variations, as shown in Fig. 2d, e. Hence, the tensile properties observed after SR are comparable to those of as-built samples. This result also shows the limited dependence of 316L tensile properties on the process-induced residual stresses.

Fatigue Properties The tension–compression (R  −1) fatigue tests performed to evaluate the impact of post-treatments on Ti–6Al–4V and on 316L are presented, respectively, in Fig. 3a, b. Figure 3a shows Ti–6Al–4V S-N curves evaluated for the three different conditions, namely, as-built, stress-relieved and HIP’ed. The graphs show that the two post-treatments, i.e. SR and HIP, improve the fatigue properties of Ti–6Al–4V samples. The poor fatigue property of as-built samples, compared with SR and HIP’ed samples, can be attributed to the presence of high residual stress levels generated during the L-PBF production process, as reported by Ramulu and Edwards [6] and to the brittle α martensitic microstructure [15].

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Fig. 4 SEM fractured surface pictures of Ti–6Al–4V samples in as-built (a, d), stress-relieved (b, e) and HIP’ed (c, f) conditions

The two heat treatments have different impacts on Ti–6Al–4V fatigue properties. On the one hand, SR produces substantial improvement of Ti–6Al–4V fatigue properties. Relief of residual stresses can be considered an influencing factor for performance enhancement after SR [5]. The other factor responsible for the improvement can be the microstructure refinement of the α martensitic microstructure into a fine lamellar α + β. In LCF fatigue regime crack initiation and propagation is influenced by α colony size in the α + β matrix. Small α colony sizes, to a certain extent, roughen the micro-crack front and hinder crack propagation, leading to improved fatigue lives [16]. Transformation from α martensitic microstructure to α + β lamella also increased the ductility of the material, which is, to some extent, beneficial for crack propagation resistance. Therefore, a lamellar α + β microstructure, relieved from residual stresses by SR, can lead to superior fatigue performance. On the other hand, HIP does not represent a source of substantial fatigue improvement. Despite relieving residual stresses, HIP coarsens the microstructure (Fig. 2c) and leads to higher ductility with a slight reduction in tensile strength. Generally, a fine-grained microstructure exhibits good crack propagation resistance, whereas a coarse-grained microstructure offers good resistance against crack initiation [5]. However, the as-built rough surface with unfused powders (Fig. 4f), possesses multiple stress raisers acting as crack initiation sites. Coarsened microstructure after HIP can lead to a reduced resistance for crack propagation and lowered fatigue performance than SR. Kasperovich et al. [5] and Chastand et al. [17] observed enhanced fatigue behavior in Ti6Al4V after HIP than SR, the reason being closing of largesized pores within the samples. Since the samples used in this research were dense and devoid of such large-sized pores, fatigue results were not overshadowed by pores but mostly influenced by microstructural and surface factors.

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The fine lamellar microstructure (Fig. 2b) observed after SR, therefore, possesses an optimum balance between strength and ductility than the other two conditions. Since it is also justified by lamellar resistance to crack initiation and propagation, the chosen SR treatment has proved to be beneficial for fatigue behavior. The analysis of Ti–6Al–4V fracture surfaces presented in Fig. 4a–c shows that the fatigue cracks always initiated from multiple points across the outer surface of the samples. Crack initiation points are associated with surface defects, such as unmolten particles, as shown in Fig. 4d–f. After the initiation, cracks propagate along the cross section of the sample until brittle failure occurs. This type of fatigue failure has been experienced for all the post-treatments investigated. It should be highlighted that during this investigation, all tested coupons failed because of surface defects, and none of the failure was due to the presence of internal defects. From this analysis, several observations can be made: (i) the surface defects can be considered the main driver for crack initiation for all investigated conditions, i.e. as-built, SR and HIP, leading L-PBF Ti–6Al–4V to under-perform compared to conventionally manufactured Ti–6Al–4V [18]; (ii) heat treatments are beneficial in terms of fatigue performance since they reduce residual stresses [6] that could negatively influence the dynamic properties; (iii) microstructure evolution from the martensitic α (Fig. 2a) to a fine lamellar α + β (Fig. 2b) can have a beneficial impact on the fatigue properties of Ti–6Al–4V; (iv) however, further HIP-induced microstructure coarsening reduces the fatigue performance compared to SR. The impact of a stress-relief post treatment on 316L fatigue performances is presented in Fig. 3b. The graph shows that the SR post-treatment slightly influences the fatigue behavior of 316L in as-built condition, leading to a small reduction in the slope of the 316L-SR S-N curve. For the two analyzed conditions, small data scatter was observed. Due to the high strength and the high ductility of AM’ed 316L, stress-raiser can be compensated; for this reason the material reaction to defects and residual stresses is relatively limited. Hence, a stress-relief operation does not induce significant changes in the S-N curves of Fig. 3b, as was also previously observed by Leuders et al. [19]. The analysis of the fractured surfaces reported in Fig. 5a, b highlights multiple crack initiations along the surface of the sample associated with surface defects generated during the manufacturing process. Similar to the observations in Ti–6Al–4V, the samples were nearly free of internal pores and defects, and cracks always initiated from the surface.

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Fig. 5 SEM fractured surface pictures of 316L samples in as-built (a–c), stress-relieved (b–d)

Conclusions The fatigue properties of Ti–6Al–4V and 316L produced by means of L-PBF were investigated on samples that received different heat treatments in as-built surface condition. Overall, the following conclusions can be drawn: 1. Stress-relief performed at 850 °C for 2 h significantly improves the fatigue properties of Ti–6Al–4V samples. 2. HIP does not improve the fatigue properties of high-quality Ti–6Al–4V parts produced by L-PBF as much as a dedicated stress-relief treatment. In this study, the low residual porosity level generated during the production was only slightly reduced by HIP. 3. The production of almost fully dense parts can have a beneficial impact on the costs associated with the post treatments of Ti–6Al–4V AM’ed parts. In order to achieve desired mechanical properties, HIP can be replaced by an economically competitive stress relieving treatment. 4. Stress-relief performed at 470 °C for 5 h did not significantly change the fatigue properties of 316L samples. 5. The samples used in this investigation were of high quality and almost pore-free test coupons; therefore, fatigue damages always occurred from surface defects

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identified as partially molten particles powder that create crack initiation favorable sites for both Ti–6Al–4V and 316L. 6. The surface defects have proven to be the main drivers for fatigue failure of high-quality L-PBF parts; therefore, further investigations are required to better understand the impact of roughness on fatigue of AM metals. Acknowledgements The work leading to this publication has been funded by the SBO project “M3-FATAM” project (HBC.2016.0446), which fits in the MacroModelMat (M3) research program, coordinated by Siemens (Siemens PLM software, Belgium) and funded by SIM (Strategic Initiative Materials in Flanders) and VLAIO (Flanders Innovation & Entrepreneurship Agency).

References 1. Kruth J-P, Levy G, Klocke F, Childs THC (2007) Consolidation phenomena in laser and powderbed based layered manufacturing. CIRP Ann Manuf Technol 56(2):730–759 2. DebRoy T et al (2018) Additive manufacturing of metallic components—process, structure and properties. Prog Mater Sci 92:112–224 3. Qiu C, Panwisawas C, Ward M, Basoalto HC, Brooks JW, Attallah MM (2015) On the role of melt flow into the surface structure and porosity development during selective laser melting. Acta Mater 96:72–79 4. Leuders S et al (2013) On the mechanical behaviour of titanium alloy TiAl6V4 manufactured by selective laser melting: fatigue resistance and crack growth performance. Int J Fatigue 48:300–307 5. Kasperovich G, Hausmann J (2015) Improvement of fatigue resistance and ductility of TiAl6V4 processed by selective laser melting. J Mater Process Technol 220:202–214 6. Edwards P, Ramulu M (2014) Fatigue performance evaluation of selective laser melted Ti–6Al–4V. Mater Sci Eng A 598:327–337 7. Cao F, Zhang T, Ryder MA, Lados DA (2018) A review of the fatigue properties of additively manufactured Ti–6Al–4V. JOM 70(3):349–357 8. Riemer A, Leuders S, Thöne M, Richard HA, Tröster T, Niendorf T (2014) On the fatigue crack growth behavior in 316L stainless steel manufactured by selective laser melting. Eng Fract Mech 120:15–25 9. Liverani E, Toschi S, Ceschini L, Fortunato A (2017) Effect of selective laser melting (SLM) process parameters on microstructure and mechanical properties of 316L austenitic stainless steel. J Mater Process Technol 249:255–263 10. Spierings AB, Wegener K, Starr TL (2013) Fatigue performance of additive manufactured metallic parts. Rapid Prototyp J 19(2):88–94 11. Zhang M, Li H, Zhang X, Hardacre D (2016) Review of the fatigue performance of stainless steel 316L parts manufactured by selective laser melting 12. Kranz J, Herzog D, Emmelmann C (2014) Design guidelines for laser additive manufacturing of lightweight structures in TiAl6V4. J Laser Appl 27(S1):S14001 13. Thijs L, Verhaeghe F, Craeghs T, Humbeeck JV, Kruth J-P (2010) A study of the microstructural evolution during selective laser melting of Ti–6Al–4V. Acta Mater 58(9):3303–3312 14. Vrancken B, Thijs L, Kruth J-P, Van Humbeeck J (2012) Heat treatment of Ti6Al4V produced by selective laser melting: microstructure and mechanical properties. J Alloy Compd 541:177–185 15. Xu W, Sun S, Elambasseril J, Liu Q, Brandt M, Qian M (2015) Ti–6Al–4V additively manufactured by selective laser melting with superior mechanical properties. JOM 67(3):668–673 16. Lütjering G (1998) Influence of processing on microstructure and mechanical properties of (α+β) titanium alloys. Mater Sci Eng A 243(1):32–45

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17. Chastand V, Tezenas A, Cadoret Y, Quaegebeur P, Maia W, Charkaluk E (2016) Fatigue characterization of titanium Ti–6Al–4V samples produced by additive manufacturing. Proc Struct Integr 2:3168–3176 18. Li P, Warner DH, Fatemi A, Phan N (2016) Critical assessment of the fatigue performance of additively manufactured Ti–6Al–4V and perspective for future research. Int J Fatigue 85:130–143 19. Leuders S, Lieneke T, Lammers S, Tröster T, Niendorf T (2014) On the fatigue properties of metals manufactured by selective laser melting—the role of ductility. J Mater Res 29(17):1911–1919

Fracture Toughness and Fatigue Strength of Selective Laser Melted Aluminium–Silicon: An Overview Leonhard Hitzler, Enes Sert, Markus Merkel, Andreas Öchsner and Ewald Werner

Abstract Metals fabricated in a powder-bed environment can achieve outstanding mechanical properties. Aluminium-silicon (AlSi) alloys can exhibit an anisotropic behaviour and exhibit inhomogeneities and predetermined sites of fracture in their asfabricated state. The work at hand provides an overview of the fracture toughness and the fatigue performance of selective laser melted AlSi, including surface treatments and the impact of the irradiation paradigm. In addition, the results of conventional post-heat treatments on selective laser melted material as well as their limitations are discussed. Keywords Additive manufacturing · Anisotropy · Inhomogeneity · Heat treatment · LCF · HCF

Introduction and Motivation Selective laser melting (SLM) is one of the most promising future production technologies for highly complex geometrical structures. Due to the layer-wise, laser beam-based fabrication method, limitations present in conventional manufacturing techniques are overcome, and because of the involved rapid cooling rates and almost instantaneous solidification of the melt, an ultra-fine-grained microstructure can be obtained [1]. The properties of selective laser melted metals are known to be L. Hitzler (B) · E. Werner Institute of Materials Science and Mechanics of Materials, Technical University Munich, 85748 Garching, Germany e-mail: [email protected] E. Sert · A. Öchsner Faculty of Mechanical Engineering, Esslingen University of Applied Sciences, 73728 Esslingen, Germany M. Merkel Faculty of Mechanical Engineering, Aalen University of Applied Sciences, 73430 Aalen, Germany © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_37

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potentially anisotropic, whereby the magnitude of anisotropy depends on the chosen raw material, and sometimes can be also inhomogeneous [2]. For SLM-aluminiumsilicon (AlSi) both is the case, whereby the inhomogeneities are twofold. On the microscale, high cooling rates lead to an increased (non-equilibrium) solubility of Si in the α-Al crystal, which consolidates in a cellular structure starting from the melt pool boundaries [3]. Correspondingly, the Si content is inhomogeneous across a single scan track, with the Si content decreasing from the scan track boundaries towards its centre. Si segregations occur in these oversaturated regions, favoured by the secondary re-melting during the fabrication of the neighbouring scan track or the subsequent layer [4, 5]. Once present, the Si-segregations hinder the Al-grains from growing further, representing an inherent grain-stabilisation feature [6–9]. This peculiarity, however, is beneficial and detrimental at the same time. On the downside, and when left untreated, these Si-enriched areas are rather brittle and cause predetermined locations for failure. Investigations revealed a noteworthy drop in the static uniaxial tensile and compressive strength, occurring at a 45° angle offset between layers and mechanical loading attributed to shear fracture along the enrichment zones [10, 11]. The inhomogeneity on the macroscale is reasoned to varying age-hardening states across a sample or component, leading to a height- and built-time dependent fluctuation in strength in the as-fabricated state [10]. Both inhomogeneities impact and limit the achievable material strength, the static and dynamic strength, and complicate the description of the material characteristics. It should be noted that both of these inhomogeneities can be reduced or even removed by means of post-heat treatments (pHT), but the homogenisation on the microscale with known standardised pHTs leads to a considerable decrease in mechanical strength [12]. The overview and discussion of results in this work focus on the few available studies on the fracture toughness and fatigue strength of selective laser melted AlSi alloys.

Discussion Fracture Toughness For AlSi10Mg, we have shown that in the as-fabricated state, its fracture toughness varied upon the angle of attack [13]. Crack propagation parallel to the layering, correspondingly in direction of the agglomerated embrittlement between consecutive layers, lowered the fracture toughness by about 20%. It should be noted that the fracture toughness does not correlate with the tensile strength. In Fig. 1, the progression of the tensile strength, i.e. yield and ultimate tensile strength, and the fracture toughness for the crack opening mode I are depicted side by side. Samples for these tests were fabricated in the same batches to completely exclude the possibilities of external influences. The anisotropy in fracture toughness of SLM-metal is material-dependent, and the weak point is not necessarily located in between consec-

Fracture Toughness and Fatigue Strength of Selective Laser …

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Fig. 1 Comparison of polar angle dependencies in AlSi10Mg: fracture toughness and tensile strength; dedicated studies on either topic can be found in [10, 13]

utive layers. Stainless steel can be highlighted as an example, showing the opposite tendency [14].

Fatigue Strength Regarding the fatigue strength, the avoidance of crack initiation points is crucial. Therefore, additional factors like inbuilt voids, such as cracks and pores, the surface condition and remaining residual stresses play an important role. Tang and Pistorius [15] showed that by reducing the hatch distance the fraction of remaining pores and oxides can be reduced, however increasing production costs and processing time. The overall bulk porosity of SLM-Al-based raw materials is rather uncritical nowadays, since almost fully dense parts can be fabricated with standardised parameter sets [1]. The critical aspect is the location of the remaining pores, which are predominantly found close to the surface, and thus, negatively impact the fatigue life [16]. Yang et al. [17] noted that the sub-surface porosity is more detrimental to the fatigue strength than the rough as-fabricated surface condition and highlighted the importance of the contour scanning schemata. Beevers et al. [18] compared samples fabricated with and without a dedicated contour irradiation and documented the best fatigue life for samples fabricated without a contour irradiation. This, of course, raises the question of why a contour irradiation is applied. Its benefit lies in the achievement of a smooth and continuous geometry with adequate surface roughness and good accuracy regarding shape and size of the component to be manufactured [1]. Neither option, however, outperforms machined samples, since cutting procedures like milling or drilling remove the sub-surface porosity, whilst improving the surface roughness as well. Surface treatments like shot-peening and electrochemical machining were

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found to be beneficial for fatigue performance, especially in the high-cycle fatigue (HCF) regime; however, due to the very gentle removal of material, any present sub-surface porosity still remains [19]. As for surface treatments without material removal, which improves the surface finish and reduces the sub-surface porosity, laser polishing might be an interesting option, as a recent study has suggested [20]. Conflicting information is found on the fatigue performance of heat-treated AlSi. This is attributed to the fact that conventional, for cast microstructure optimised heat treatments (only considering pHTs which include a solution annealing step) lead to both positive and negative effects on the fatigue performance, depending on the considered aspect. PHTs, like the very common T6 treatment, lead to a relaxation of residual stresses, the removal of inhomogeneities (on the micro- and macroscale), accompanied by the corresponding obliteration of the anisotropy and a reduction in the number of crack initiation sites, whilst improving the ductility of the material, but at the cost of a reduction in overall strength [6, 12, 18, 21–23]. In summary, the following conclusion of pHTs on the fatigue performance can be drawn: For low-cycle fatigue (LCF) the loss in overall strength due to a coarser microstructure shows reduced fatigue performance, whereas for HCF the studies mostly document beneficial results.

Perspective and Outlook At the moment, more in-depth studies on the fracture toughness of as-fabricated AlSi are undertaken, to identify the critical angle and/or to explain the 20% loss in fracture toughness. As for pHTs and their impact on the fracture behavior, a similar picture is anticipated, whereby the achieved homogenization will eliminate the directional dependency, but lower the fracture toughness altogether. The major conclusion in either case is that at the current state, an appropriate pHT, optimised for the SLMmicrostructure, needs to be developed. In view of this, the current and following steps include—besides more in-depth studies of the irradiation schemata on as-fabricated parts—the development of a dedicated heat treatment which ideally combines the homogenisation and the corresponding improvements of the properties across all areas by dissolving the inbuilt embrittlement, without causing a major coarsening of the microstructure. One promising approach to achieve this goal is seen in performing heat treatments at elevated pressures, for example, in combination with hot isostatic pressing (HIP) [24, 25]. This enables to influence the diffusivity of the material and has the potential to alleviate the grain coarsening problematic, whilst gaining the benefit of a further reduction in the number of crack initiation sites.

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References 1. Hitzler L, Merkel M, Hall W, Öchsner A (2018) A review of metal fabricated with laserand powder-bed based additive manufacturing techniques: process, nomenclature, materials, achievable properties, and its utilization in the medical sector. Adv Eng Mater 20:1700658. https://doi.org/10.1002/adem.201700658 2. Hitzler L, Hirsch J, Heine B, Merkel M, Hall W, Öchsner A (2017) On the anisotropic mechanical properties of selective laser melted stainless steel. Materials 10:1136. https://doi.org/10. 3390/ma10101136 3. Prashanth KG, Scudino S, Klauss HJ, Surreddi KB, Löber L, Wang Z, Chaubey AK, Kühn U, Eckert J (2014) Microstructure and mechanical properties of Al–12Si produced by selective laser melting: effect of heat treatment. Mater Sci Eng A 590:153–160. https://doi.org/10.1016/ j.msea.2013.10.023 4. Aboulkhair NT, Tuck C, Ashcroft I, Maskery I, Everitt NM (2015) On the precipitation hardening of selective laser melted AlSi10Mg. Metall Mater Trans A 46:3337–3341. https://doi. org/10.1007/s11661-015-2980-7 5. Tang M, Pistorius PC (2017) Anisotropic mechanical behavior of AlSi10Mg parts produced by selective laser melting. JOM 69:516–522. https://doi.org/10.1007/s11837-016-2230-5 6. Aboulkhair NT, Maskery I, Tuck C, Ashcroft I, Everitt NM (2016) Improving the fatigue behaviour of a selectively laser melted aluminium alloy: Influence of heat treatment and surface quality. Mater Des 104:174–182. https://doi.org/10.1016/j.matdes.2016.05.041 7. Takata N, Kodaira H, Sekizawa K, Suzuki A, Kobashi M (2017) Change in microstructure of selectively laser melted AlSi10Mg alloy with heat treatments. Mater Sci Eng A 704:218–228. https://doi.org/10.1016/j.msea.2017.08.029 8. Tang M (2017) Inclusions, porosity, and fatigue of AlSi10Mg parts produced by selective laser melting. Ph.D. thesis, Carnegie Mellon University 9. Ding Y, Muñiz-Lerma JA, Trask M, Chou S, Walker A, Brochu M (2016) Microstructure and mechanical property considerations in additive manufacturing of aluminum alloys. MRS Bull 41:745–751. https://doi.org/10.1557/mrs.2016.214 10. Hitzler L, Janousch C, Schanz J, Merkel M, Heine B, Mack F, Hall W, Öchsner A (2017) Direction and location dependency of selective laser melted AlSi10Mg specimens. J Mater Process Tech 243:48–61. https://doi.org/10.1016/j.jmatprotec.2016.11.029 11. Hitzler L, Schoch N, Heine B, Merkel M, Hall W, Öchsner A (2018) Compressive behaviour of additively manufactured AlSi10Mg. Mat -wiss u Werkstofftech 49:683–688. https://doi.org/ 10.1002/mawe.201700239 12. Aboulkhair NT (2016) Additive manufacture of an aluminium alloy: processing, microstructure, and mechanical properties. Ph.D. thesis, University of Nottingham 13. Hitzler L, Hirsch J, Schanz J, Heine B, Merkel M, Hall W, Öchsner A (2017) Fracture toughness of selective laser melted AlSi10Mg. P I Mech Eng L J Mat. https://doi.org/10.1177/ 1464420716687337 (Online first) 14. Riemer A, Leuders S, Thöne M, Richard HA, Tröster T, Niendorf T (2014) On the fatigue crack growth behavior in 316L stainless steel manufactured by selective laser melting. Eng Fract Mech 120:15–25. https://doi.org/10.1016/j.engfracmech.2014.03.008 15. Tang M, Pistorius PC (2017) Oxides, porosity and fatigue performance of AlSi10Mg parts produced by selective laser melting. Int J Fatigue 94:192–201. https://doi.org/10.1016/j.ijfatigue. 2016.06.002 16. Hitzler L, Janousch C, Schanz J, Merkel M, Mack F, Öchsner A (2016) Non-destructive evaluation of AlSi10Mg prismatic samples generated by selective laser melting: influence of manufacturing conditions. Mat -wiss u Werkstofftech 47:564–581. https://doi.org/10.1002/mawe. 201600532 17. Yang KV, Rometsch P, Jarvis T, Rao J, Cao S, Davies C, Wu X (2018) Porosity formation mechanisms and fatigue response in Al-Si-Mg alloys made by selective laser melting. Mater Sci Eng A 712:166–174. https://doi.org/10.1016/j.msea.2017.11.078

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18. Beevers E, Brandão AD, Gumpinger J, Gschweitl M, Seyfert C, Hofbauer P, Rohr T, Ghidini T (2018) Fatigue properties and material characteristics of additively manufactured AlSi10Mg—effect of the contour parameter on the microstructure, density, residual stress, roughness and mechanical properties. Int J Fatigue 117:148–162. https://doi.org/10.1016/j. ijfatigue.2018.08.023 19. Uzan NE, Ramati S, Shneck R, Frage N, Yeheskel O (2018) On the effect of shot-peening on fatigue resistance of AlSi10Mg specimens fabricated by additive manufacturing using selective laser melting (AM-SLM). Addit Manuf 21:458–464. https://doi.org/10.1016/j.addma.2018.03. 030 20. Schanz J, Hofele M, Ruck S, Schubert T, Hitzler L, Schneider G, Merkel M, Riegel H (2017) Metallurgical investigations of laser remelted additively manufactured AlSi10Mg parts. Mat -wiss u Werkstofftech 48:463–476. https://doi.org/10.1002/mawe.201700039 21. Buchbinder D, Meiners W, Brandl E, Palm F, Müller-Lohmeier K, Wolter M, Over C, Moll W, Weber J, Skrynecki N, Grad J, Neubert V (2010) Abschlussbericht - Generative Fertigung von Aluminiumbauteilen für die Serienproduktion, 01RIO639A-D, BMBF, Fraunhofer ILT 22. Brandl E, Heckenberger U, Holzinger V, Buchbinder D (2012) Additive manufactured AlSi10Mg samples using selective laser melting (SLM): microstructure, high cycle fatigue, and fracture behavior. Mater Des 34:159–169. https://doi.org/10.1016/j.matdes.2011.07.067 23. Maskery I, Aboulkhair NT, Tuck C, Wildman RD, Ashcroft IA, Everitt NM, Hague RJM (2015) Fatigue performance enhancement of selectively laser melted aluminium alloy by heat treatment. Paper presented at SFF symposium, Austin, Texas, USA 24. Hafenstein S, Brummer M, Ahlfors M, Werner E (2016) Combined hot isostatic pressing and heat treatment of aluminum A356 cast alloys. HTM J Heat Treat Mater 71:117–124. https:// doi.org/10.3139/105.110281 25. Hafenstein S, Brummer M, Ahlfors M, Werner E (2016) Kombiniertes Heißisostatisches Pressen (HIP) und Wärmebehandlung von einer A356 Aluminiumgusslegierung. Druckguss 7–8:316–321

The Effect of Heat Treatment and Alloying of Ni–Ti Alloy with Copper on Improving Its Fatigue Life Wisam Abu Jadayil and Duaa Serhan

Abstract Ni–Ti alloys have achieved great importance in industry, mainly for their innovative use in practical medical applications. A major reason for that is their fatigue lives. In this research, fatigue life of different Ni–Ti alloys has been investigated for different compositions of Ni–Ti alloy samples and different percentages of copper additions. Three major compositions were investigated under different fatigue loadings. The first group of samples had a composition of Ni47.8Ti42.2, and the remaining 10% was copper. A second group of samples had a composition of Ni52.8Ti42.2 and 5% of copper, and a third group pf samples had a composition of Ni47.8Ti47.2 and 5% of copper. These samples were prepared by casting. Comparison between fatigue lives of the prepared samples and the Ni52.8Ti47.2 sample was made, once without heat treatment of samples and once with heat treatment to find the best alloy composition with the best fatigue life. It was found that adding copper would improve the fatigue life of NiTi such that Ni has significantly higher percentages than Ti. Moreover, heat treating the NiTiCu alloy would improve its fatigue life by almost 10%. Keywords Ni–Ti alloy · Fatigue life · Copper addition · Heat treatment

Introduction and Literature Review The first nitinol alloy was prepared by Buhler in the Naval Ordinance Laboratory in 1962 [1, 2]. On the other hand, the first fatigue study of NiTi SMAs was performed by Melton and Mercier [3] in 1978, where pseudoelastic fatigue tests were run on wire specimens with different temperatures. That work was soon followed by the W. Abu Jadayil (B) Mechanical and Industrial Engineering, American University of Ras Al Khaimah, Ras Al Khaimah 10021, UAE e-mail: [email protected] D. Serhan Industrial and Systems Engineering, State University of New York at Binghamton, Binghamton, NY 13902, USA e-mail: [email protected] © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_38

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work of McNichols and Brookes [4], who studied the fatigue life of NiTi springs. It was not until the early 90s, when the medical industry began to push for less invasive medical procedures and alternative implants [5]. Takeshita et al. [6] implanted cylindrical NiTi parts in rats for 168 days in the year 1997. An electropolished nitinol sample was implanted in a periosteum osteoblasts; for the first 26th week, it showed no toxicity effects but a small deceleration in proliferation process [7]. Nickel percentage of nitinol composition where the issue that studied by Kapanen et al. [8] in 2002. A fabricated sample with different percentages of nickel and titanium had been investigated for in vitro studies by Bogdanski et al. [9]. The highest biological compatibility assured to be within a maximum 50% nickel element of the alloy weight, higher percentages of Ni have revealed nickel releases and the released nickel rapidly reached cytotoxic concentrations within 1 day. Many similar researches later handled toxicity and corrosion resistance of the NiTi alloy [10–16]. It was found that the “memorial” effect point of transitions from structure to another fits perfectly body temperature. Equiatomic nitinol, with its pseudoelastic effect, was found to have several ideal properties for such aim. Fortunately, Mcklevey and Ritchie [17] found that (Ni50Ti50) casted alloy samples showed a full Austenitic structure at body temperature, which means a perfect mechanical condition providing super-elasticity. Es-Souni et al. [18] investigated the mechanical behavior for 55.8 wt% Ni–Ti samples. Thermo-mechanical stress–strain relationships and deformation processes had been investigated in detail by Sittner et al. [19]. Studying the mechanical properties, especially the fatigue and tensile properties, of alloys was the main focus of many researchers [20–35]. Abu Jadayil and Alnaber [24] found that the composition of the Ni–Ti alloy that gives the best fatigue life is Ni52.8Ti47.2. That was the start point of this research, where copper of percentages 5 and 10% were added to the Ni–Ti alloy, by reducing 5% from Ni and 5% from Ti, reducing only 5% from Ti and a third group of samples where the 5% was reduced from Ni.

Methodology The methodology followed in this research went through three major stages; sample preparation with the specified percentages of nickel, titanium and copper, then the heat treatment stage of 50% of the samples, and the final stage was testing all heattreated and non-heat-treated samples by fatigue cyclic loading.

Sample Preparation Nickel, titanium, copper and tin metals were used each with the size and purity of powders as shown in Table 1.

The Effect of Heat Treatment and Alloying of Ni–Ti Alloy … Table 1 Specifications of used metals

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Metal

Purity (%)

Melting temp. (c)

Shape

Ni

99.8

1453

Plates and bullets

Ti

99.6

1670

Powder (150 μm)

Cu

99.2

1083

Powder

Fig. 1 Fatigue life testing sample

The three components were converted to molten, mixed with the required percentages and then poured in a mold that produced fatigue samples as the one shown in Fig. 1. The first group (G1) of prepared samples had a composition of Ni47.8Ti42.2, and the remaining 10% was copper. A second group (G2) of prepared fatigue samples had a composition of Ni52.8Ti42.2 and 5% of copper and a third group (G3) of prepared samples had a composition of Ni47.8Ti47.2 and 5% of copper. One more fourth group (G4) of samples was made with a composition of Ni52.8Ti47.2 where no copper was added.

Heat Treatment Each of the four groups of samples was divided into two types; one heated treated and the other is not heat-treated. The heat-treated samples were heated to a temperature of 400 °C, kept at that temperature for 2 h, and then cooled at air to room temperature. The 400 °C was chosen because it is very close to recrystallization temperature of the copper.

Fatigue Life Testing Samples from the four groups, heat-treated and without heat treatment, were tested on fatigue life testing machine under fatigue loadings of 28 N, 39 N, 51 N, 74 N and 92 N, with a loading frequency of 60 Hz, and average behavior of the samples was recorded as shown in Table 2. Figure 2 shows one of the Ni52.8Ti47.2 samples after failure. The stress amplitude was calculated according to the following equation: σa  32LF/πd3

(1)

Stress (MPa)

125

175

230

330

410

Load (N)

28

39

51

74

92

30,568

60,348

271,381

4,120,682

579,270,163

G4

32,753

65,236

301,341

4,512,146

648,782,587

G4h

34,542

68,190

306,669

4,656,332

654,575,301

G1

37,993

75,025

337,384

5,121,997

720,032,903

G1h

Table 2 Fatigue life (in cycles) of prepared samples under different fatigue loadings

40,349

79,638

358,334

5,439,401

764,637,341

G2

43,983

86,855

390,702

5,929,118

833,454,582

G2h

25,831

50,044

225,311

3,420,512

480,794,702

G3

27,921

54,976

246,912

3,749,840

527,135,876

G3h

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Fig. 2 Fatigue sample after failure

where F is the applied weight, L is the distance between applied force point and the center of the tested sample, which was 28 mm, and d is the diameter of the sample at its center, which was 4 mm.

Discussion of the Results Results shown in Table 2 were plotted for the four groups, as shown in Fig. 2. Except the G2 and G2h (where h refers to heat-treated sample), all other group samples showed improved fatigue lives compared to the G4 samples where no copper was added (Fig. 3). Comparing G4 with G4h samples shows that heat-treated samples can live on average from 7 to 12% more. G1 samples live on average around 13% more than G4 samples. That proves adding copper with 10% of weight resulted in improving the fatigue life. In G1 samples, the added 10% of copper was deducted equally from Ti and Ni, 5% from each. As expected, G1h samples showed around 10% improvement of fatigue lives compared to G1 samples, and around 25% improvement in fatigue life compared to the basic samples of G4. The best fatigue life achieved was using G2 and G2h samples, where copper was added with 5%, which was deducted from Ti%. G2 samples showed improvements in fatigue life of 32% on average compared to G4 samples. When heat-treated, G2h samples showed improvements in fatigue life of 9% compared to G2 samples without heat treatment, and 44% compared to basic samples of G4. G3 and G3h samples showed exceptional behavior, where the fatigue life was reduced compared to G4 samples, by 17 and 9%, respectively. Although 5% copper was added to G3 samples, that percentage was taken completely from the Ni%, which resulted in having equal percentages of Ni and Ti. Having the samples of G3h heat-treated improved their fatigue life by a percentage of 9% compared to G3 samples, but they still living 9% less than the basic samples of G4.

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G4 G4h G1 G1h G2 G2h G3 G3h

410

350 330

Stress (MPa)

300 250 230

200 175

150 125

100 50 0 10000

100000

1000000

10000000

100000000

1E+09

Cycles (Rev) Fig. 3 The fatigue stress versus the failure cycles for the four tested groups

Conclusions The fatigue life of TiNi was studied by adding copper and heat-treating the samples at 400 °C for 2 h. Generally, adding copper to NiTi alloy will improve its fatigue life, such that the percentage of Ni is still higher than that of Ti. Compared to Ni52.8Ti47.2 samples, Ni52.8Ti42.2Cu5.0 samples showed the most improved fatigue life, then Ni47.8Ti42.2Cu10.0 samples. On the other hand, it was found that adding copper and making the percentages of Ti and Ni almost equal would reduce the fatigue life, as it was found when testing Ni47.8Ti47.2Cu5.0 sample. So, to improve the fatigue life of NiTi alloy, Cu should be added by almost 5% and the Ni% should be kept higher than Ti%. In all cases, heat-treating the samples would improve their fatigue life by almost 10%.

References 1. Buehler WJ, Gilfrich JV, Wiley RC (1963) Effect of low-temperature phase changes on the mechanical properties of alloys near the composition of TiNi. J Appl Phys 34:1475–1477 2. Buehler WJ, Wang FE (1967) A summary of recent research on the nitinol alloys and their potential application in ocean engineering. Ocean Eng J 1:105–120

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3. Melton KN, Mercier O (1978) Fatigue of niti thermoelastic martensites. Acta Metall 27:137–144 4. McNichols JL, Brookes PC, Cory JS (1981) NiTi fatigue behavior. J Appl Phys 52:7442–7444 5. Chengli S (2010) History and current situation of shape memory alloys devices for minimally invasive surgery. Open Med Devices J 2:24–31 6. Takeshita F, Takata H, Ayukawa Y, Suetsugu T (1997) Histomorphometric analysis of the response of rat tibiae to shape memory alloy (nitinol). Biomater J 18:21–25 7. Shabalovskaya SA (2001) Physicochemical and biological aspects of nitinol as a biomaterial. Int Mater J 46:1–18 8. Kapanen A, Ilvesaro J, Danilov A, Ryhänen J, Lehenkari P, Tuukkanen J (2002) Behaviour of nitinol in osteoblast-like ROS-17 cell cultures. Biomater J 23:645–650 9. Bogdanski D, Köller M, Müller D, Muhr G, Bram M, Buchkremer HP, Stöver D, Choi J, Epple M (2002) Easy assessment of the biocompatibility of Ni–Ti alloys by in vitro cell culture experiments on a functionally graded Ni–NiTi–Ti material. Biomater J 23:4549–4555 10. Wu S, Liu X, Chan YL, Ho JP, Chung CY (2007) Nickel release behavior, cytocompatibility, and superelasticity of oxidized porous singlephase NiTi. J Biomed Mater 81:948–955 11. Liu XM, Wu SL, Chan YL, Chu PK, Chung CY (2007) Surface characteristics, biocompatibility, and mechanical properties of nickel–titanium plasma implanted with nitrogen at different implantation voltages. J Biomed Mater 82:469–478 12. Rhalmi S, Charette S, Assad M, Coillard C, Rivard C (2007) The spinal cord dura mater reaction to nitinol and titanium alloy particles: a 1-year study in rabbits. Eur Spine J 16:1063–1072 13. Shishkovsky I, Yu M, Smurov I (2007) Nanofractal surface structure under laser sintering of titanium and nitinol for bone tissue engineering. J Appl Surf Sci 254:1145–1149 14. Rocher P, El Medawar L, Hornez C, Traisnel M, Breme J, Hildebrand H (2004) Biocorrosion and cytocompatibility assessment of NiTi shape memory alloys. Scr Mater J 50:255–260 15. Kujala S, Tuukkanen J, Jamsa J, Danilov A, Paramila A (2002) Comparison of bone modeling effects caused by curved and straight Nickel–Titanium nails. J Mater Sci 13:1157–1161 16. Sun EX, Fine S, Nowak WB (2002) Electrochemical behavior of nitinol alloy in ringer‘s solution. J Mater Sci 13:959–964 17. McKelvey AL, Ritchie RO (2001) Fatigue-crack growth behavior in the superelastic and shapememory alloy nitinol. Metall Mater Trans 32A:731–743 18. Es-Souni M, Es-Souni M, Brandies H (2001) On the transformation behaviour, mechanical properties and biocompatibility of two NiTi-based shape memory alloys: NiTi42 and NiTi42Cu7. J Biomater 22:2153–2216 19. Sittner P, Landa M, Lukas P, Novak V (2006) R-phase transformation phenomena in thermomechanically loaded NiTi polycrystals. Mech Mater J 38:475–492 20. Khraisat W, Abu Jadayil W, Al-Zain Y, Mismar S (2018) The effect of rolling direction on Welded DP 1000 steel microstructure and its mechanical properties. Cogent Eng 5(1):1–11 21. Khraisat W, Abu Jadayil W, Rawashdeh N, Borgstrom H (2018) The role of phosphorus in pore rounding of sintered steels. Cogent Eng 5(1):1–12 22. Abu Jadayil W (2015) Surface and subsurface defects investigation of Ni-Ti samples processed by different fabrication methods. Res J Appl Sci Eng Technol 11(11):1190–1195 23. Abu Jadayil W, Alnaber M (2015) Experimental investigation of tensile properties of Ni-Ti samples prepared by different techniques. Int J Appl Eng Res 10(6):15651–15659 24. Abu Jadayil W, Alnaber M (2014) Assessment of fatigue life of Ni-Ti samples prepared by different techniques. Appl Mech Mater 477–478:1264–1268 25. Abu Jadayil W, Mohsen M (2011) Experimental investigation of self actuating traction drives with solid and hollow rollers. Int Rev Mech Eng 5(4):637–645 26. Abu Jadayil W (2010) Revision of the recent heterogeneous object modeling techniques. Jordan J Mech Ind Eng 4(6):779–788 27. Abu Jadayil W, Mohsen M (2010) Design and manufacturing of self-actuating traction drives with solid and hollow rollers. Jordan J Mech Ind Eng 4(4):467–476 28. Khraisat W, Abu Jadayil W (2010) Strengthening aluminum scrap by alloying with iron. Jordan J Mech Ind Eng 4(3):372–377

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29. Abu Jadayil W, Jaber N (2010) Numerical prediction of optimum hollowness and material of hollow rollers under combined loading. J Mater Design 31(3):1490–1496 30. Abu Jadayil W, Khraisat W (2010) Predicting the optimum hollowness of normally loaded cylindrical rollers using finite element analysis. J Mater Sci Technol 26(2):176–183 31. Khraisat W, Borgstrom H, Nyborg L, Abu Jadayil W (2009) Optimising grey iron powder compacts. J Powder Metall 52(4):291–297 32. Abu Jadayil W (2008) Relative fatigue life estimation of cylindrical hollow rollers in general pure rolling contact. J Tribotest 14:27–42 (Wiley InterScience) 33. Abu Jadayil W, Flugrad D (2007) Fatigue life investigation of solid and hollow rollers under pure normal loading. J Tribotest 13:165–181 (Wiley InterScience) 34. Abu Jadayil W (2011) experimental investigation of solidification time effects on surface and subsurface aluminum casting defects. Int Rev Mech Eng 5(4):569–576 35. Abu Jadayil W (2011) Studying the effects of varying the pouring rate on the casting defects using non-destructive testing techniques. Jordan J Mech Ind Eng 5(6):521–526

Effect of Adding Yttrium on the Inclusion Modification and Impact Toughness of E36 Shipbuilding Steel Xiaojun Xi, Maolin Ye, Shufeng Yang and Jingshe Li

Abstract The scientific exploration and exploitation of the ocean resource is now proceeding at a greatly accelerated rate. Consequently, it is required that shipbuilding steel must possess high strength and toughness. In this comparative study, the inclusion, microstructure and impact toughness of an E36 shipbuilding steel, with and without addition of yttrium, were investigated. The results show that the elongated MnS inclusions in E36 steel were replaced by spindle and spherical inclusions containing yttrium upon addition of 0.023 wt% yttrium, leading to the formation of the E36Y steel. The microstructure of test steels was characterized through the ferrite and pearlite phases. The addition of yttrium decreased the pearlite lamellar spacing and refined the pearlite laminae. Furthermore, the impact toughness of test steel increased significantly at different temperatures; both the longitudinal and transverse impact fracture displayed ductility characteristics; and the anisotropy of longitudinal and transverse impact toughness decreased significantly. Keywords Yttrium · E36 shipbuilding steel · Inclusion modification · Impact toughness

Introduction The scientific exploration and exploitation of the ocean resource is now proceeding at a greatly accelerated rate [1]. Consequently, it is required that shipbuilding steel must possess high strength and toughness, excellent corrosion resistance and high heat input welding performance. However, non-metallic inclusions, generated inevitably due to active elements in steel contacting and reacting with refractory materials and slags during steelmaking and casting process, are harmful to steel properties such as strength, toughness, fatigue, and corrosion resistance [2], which X. Xi · M. Ye · S. Yang (B) · J. Li School of Metallurgical and Ecological Engineering, University of Science and Technology Beijing, Beijing 100083, China e-mail: [email protected] © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_39

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should be removed or modified as much as possible so it did not deteriorate steel properties. Many researchers have investigated improving the behavior of non-metallic inclusions during solidification of steel. It was found that depending on the chemical composition and size of inclusion, they can have different impacts on the mechanical properties of steel [3]. Inclusions larger than 10 µm are probable to lower the strength and toughness obviously. However, the strength would increase remarkably for steels with inclusions less than 0.3 µm [4]. With improved understanding of fine inclusions, particularly their positive influences on steel microstructures. Adding a or several elements into steel has been attempted by steel manufactures in order to form finer grain structures in steel and to improve the performance of the steel [4]. Titanium, aluminum and magnesium have a strong affinity with oxygen, and their functions have been researched a lot to improve the mechanical properties of steel [5–8]. However, as previous research has reported, Ti, Mg and Al cannot desulfurize to a very low level [9] and their ability of deoxidization is limited [4]. In addition, Al forms large Al2 O3 inclusion that would cluster [10], and the technique to add Mg in liquid steel is not mature to obtain stable Mg-added steel. Rare earth (RE) element has very strong chemical activity because of their unique electronic structures, where the valence state of the 4f channel is variable [11]. RE elements have a strong affinity to O and S, and could react with O and S with the result of forming the high-melting point REx Oy , REx Sy , REx Oy Sz , spheroidizing inclusions (such as MnS) to avoid the anisotropy of mechanical properties in final rolling products [12]. In addition, RE can dissolve into Fe with a solubility is around 10−6 ~ 10−5 ppm magnitude, and playing the role of micro-alloying [11]. Tiny amounts of an RE metal dissolved in steel can distort the iron crystal lattice and enhance the mechanical properties. RE metals tend to segregate at grain boundaries and eliminate the local weaknesses due to sulfur and phosphor atoms in steel, improving the strength of grain boundaries and shock resistance [13]. The key factors for the nucleation of intergranular bainite or acicular ferrite are the control of austenite grain size, which are both considered as favorable phases for mechanical properties at room temperature [12]. As reported [14], in carbon RE steel, RE atoms tend to segregate at the interface of ferrite and cementite due to their large radius and high aberration energy. RE atoms are thus mainly distributed at the interface of cementite alloys and grain boundaries. The grain sizes of austenite decrease significantly with the addition of a greater amount of RE metals, thus prompting the AF nucleation. In recent years, extensive research has been carried out on the influence of RE elements lanthanum and cerium on microstructures of various test steels [13–15]. Yttrium (Y), another reactive element, not only exhibits some similarities with lanthanum and cerium but also displays its own favorable characteristics [16]. Unfortunately, to the best of the authors’ knowledge, few studies have considered yttriummodified steel, especially in terms of the microstructural changes and associated influence on mechanical properties. In this paper, the inclusion, microstructure and impact toughness of E36 and E36Y shipbuilding steel were compared. The mechanism of Y affects the E36Y steel was also studied.

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Table 1 Chemical compositions of E36 and E36Y steel plate (wt%) Element

C

E36

0.088 0.118 1.456 0.027 0.006 0.0028 0.052 0.016 0.005 0.044 0.036

Si

Mn

P

S

O

Cr

Ni

Mo

Cu

Al

Y

E36Y

0.100 0.240 1.370 0.018 0.004 0.0015 0.063 0.016 0.007 0.039 0.031 0.023

Experimental Design Material Preparation The experimental E36 shipbuilding steels were prepared by remelting the commercial E36 steels in a vacuum induction melting furnace (25 kg). The E36Y steels were prepared by a vacuum induction melting furnace, the commercial E36 steels were used as the starting material, the yttrium ferroalloy was subsequently added. Here, the yttrium ferroalloy is composed of 65 wt%Y and 35 wt%Fe in this study. The final compositions of E36 and E36Y shipbuilding steels were shown in Table 1. The as-made E36 and E36Y steels were forged into 90 mm × 90 mm ingots, and hot rolled into thick plates of 29 mm in thickness with a rolling temperature at 1363 K (1090 °C) and finished at 1085 K (812 °C). Finally, the plates were air cooled to room temperature.

Methods of Analysis Detection of Inclusions In order to clarify the inclusions in steel samples, 10 × 10 × 10 mm3 cubic samples were cut from sample steel, then ground and polished using 3.5 and 0.5 µm diamond compound. The types of inclusions and their morphology were extensively analyzed by scanning electron microscope (SEM) and energy-dispersive spectrometry (EDS) (model ProX).

Microstructure Examination After etching the polished surface of steel samples with 4 pct Nital for 50 s, the steel microstructure was examined with the aid of the SEM (model Quanta).

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Fig. 1 Sampling diagram of impact test

Mechanical Properties Tests Charpy v-notch impact test specimens with dimensions of 10 mm × 10 mm × 55 mm were prepared by spark cutting from the plates and then mechanically polished. Figure 1 shows the sampling position at a quarter of the width of the steel plates, and the longitudinal specimens were cut along axes parallel to the rolling direction, while the transverse specimens are perpendicular to the rolling direction. The impact toughness tests were conducted according to ASTM E23, with a JBDW-300D testing machine at 20, 0, −20, −40, and −60 °C, respectively. All the measurements were repeated at least three times. The fracture surface of impact specimens was observed using SEM.

Results and Discussion Effect of Y on the Morphology and Type of Inclusions Figure 2 shows the morphology and type of inclusions in E36 and E36Y steel, respectively. Without adding rare earth Y, MnS is the dominating inclusion in E36 steel, which is presented in Fig. 2a. The strip-like MnS inclusion is distributed along the rolling direction and the size reaches 30 µm. By adding Y (E36Y), the strip-like sulfide inclusions become rare earth complex inclusions with spindle and spheric in shapes, and the size lower than 5 µm, as shown in Fig. 2b. Compared to Mn and Fe, Y has a stronger affinity towards O and S and is more prone to react with them to form oxide and sulphide complex inclusions. The experimental results are in accordance with the thermodynamic calculations.

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Fig. 2 SEM images and EDS of inclusions in E36 and E36Y steel: a E36 steel; b and c E36Y steel

Effect of Y on the Microstructure The microstructure of E36 and E36Y steels is shown in Fig. 3. Figure 3a, b is the longitudinal specimens, which were cut along axes parallel to the rolling direction; Fig. 3c, d is the transverse specimens, which were cut perpendicularly to the rolling direction. It can be seen that the microstructure consists entirely of ferrite and pearlite, which in turn is a mixed structure of ferrite and cementite. Compared with the microstructure of the E36 steel, shown in Fig. 3a, c, the pearlite lamellar spacing in the E36Y steel is lower and the pearlite laminae are finer, as shown in Fig. 3b, d. In E36Y shipbuilding steel, Y atoms can only replace the Fe atoms of cementite or exist as unstable Y carbides [17]. Y atoms tend to segregate on the grain boundaries of ferrite and cementite, as the lattice distortion energy carried out by Y atoms dissolving in the ferrite is much higher than that in the grain boundaries.

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Fig. 3 Microstructure and morphology of E36 and E36Y steel: a longitudinal of E36 steel, b longitudinal of E36Y steel, c transverse of E36 steel, and d transverse of E36Y steel

Therefore, Y atoms predominantly distribute in the interface and alloy cementite [18]. During the pearlite transformation, the diffusion velocity of Y atoms is smaller than carbon atoms; thus, the nucleation rate and coarsening rate of cementite depend on the diffusion and enrichment of Y atoms. Consequently, Y can prolong the incubation time of pearlite transformation and increase condensate depression of pearlite transformation, which will decrease the pearlite lamellar spacing and refine the pearlite laminae [18]. The relationship of the diffusion coefficient Di of i element and interaction coefj ficient ei of i element and j element is given by Eq. (1). j

Di  D0 exp(ei x j )

(1) j

where Di is the diffusion coefficient of i element; ei is the interaction coefficient of i element and j element; D0 is the diffusion coefficient of i element without j element; and x j is the molar fraction of j element. Y can decrease the diffusion coefficient of C and control the precipitation of carbide, which will prolong the pearlite transformation and refine the pearlite laminae.

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Fig. 4 Longitudinal and transverse impact toughness of E36 and E36Y steels

Fig. 5 Ratio of longitudinal and transverse impact toughness of E36 and E36Y steels

Effect of Y on the Impact Toughness Figure 4 exhibits the effect of Y on the impact toughness of E36 and E36Y shipbuilding steel at different test temperatures. By adding the Y, the longitudinal and transverse impact toughness increase significantly. The impact toughness of E36 steel begins to decrease below 0 °C and decrease greatly below −40 °C, while the impact toughness of E36Y steel decreases gradually at room temperature and low temperature. In addition, the ratio of longitudinal to transverse impact toughness is lower for E36Y compared with E36 steel, as shown in Fig. 5. The results indicate that the mechanical anisotropy of the E36 shipbuilding steel is reduced by adding the Y. The representative impact fractography of both steels is shown in Fig. 6. Figure 6a, b is the longitudinal impact fracture, and Fig. 6c, d is the transverse impact fracture. Both the longitudinal impact fractures display the characteristic

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Fig. 6 SEM fractography of E36 and E36Y steels: a longitudinal of E36 steel, b longitudinal of E36Y steel, c transverse of E36 steel, and d transverse of E36Y steel

of ductile. Compared with E36 steel, the fracture surface in E36Y steel is much rougher and presents more dimples, which proved the impact toughness of E36Y steel is superior to the E36 steel. For the transverse impact fracture, the E36 steel displays the characteristic of cleavage, while the E36Y steel presents the ductile fracture, and there are many cracked inclusions exist in the dimples. Therefore, the transverse impact toughness of E36 steel was improved by adding the Y. Non-metallic inclusions were an important factor influencing the fracture behavior of E36 and E36Y steels. The strip-like MnS inclusion in E36 steel distributes along the rolling direction and deteriorates the continuity of steel matrix. In addition, due to the thermal expansion coefficient and plastic deformation capacity between MnS inclusion and steel matrix have much different, there generates a stress concentration at the top of MnS inclusion in the process of rolling and heat treatment [13], thus leading to the MnS inclusion acts as the nuclei of cleavage fracture of transverse impact section. Therefore, the synthetic action deteriorates the transverse impact toughness greatly, while the longitudinal impact toughness is less affected. In E36Y steel, elongated MnS inclusion was replaced by spindle and spherical inclusio ns containing Y, and the inclusions and steel matrix have similar thermal expansion coefficient, thus the stress concentrations are avoided during rolling and heat treatment [13, 17]. Therefore, the transverse impact toughness of E36Y steel was improved greatly, comparing with the E36 steel, the impact anisotropy was optimized significantly.

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It is obvious that the addition of the Y has an important effect in inhibiting the pearlite transformation, thus decreasing the pearlite lamellar spacing and refining the pearlite laminae, the volume fraction of pearlite decreased. Generally speaking, the impact toughness of the steel would get increased with the reduction of the pearlite. These proved the longitudinal and transverse impact toughness were improved by adding the Y.

Conclusions (1) The strip-like MnS inclusions were modified into spindle and spherical inclusions containing Y by adding the Y, and the size decreases from 30 µm to less than 5 µm. (2) In E36 and E36Y steels, the microstructure consists entirely of ferrite and pearlite. The pearlite lamellar spacing decreased and the pearlite laminae were refined by following the addition of the Y. (3) By adding the Y, the longitudinal and transverse impact toughness increased significantly at different temperatures, and the impact anisotropy of longitudinal and transverse specimens decreased. All the longitudinal impact fractures displayed the characteristics of ductility. With regard to the transverse impact fracture, E36 steel displays the cleavage characteristic of fragile fracture, while E36Y steel exhibits the ductile fracture, and there are many cracked inclusions exist in the dimples. Acknowledgements The authors gratefully acknowledge the support by the National Natural Science Foundation of China (NSFC, Nos. 51664021 and 51474085), and Key Project of Natural Science Foundation of Jiangxi Province (No. 20171ACB20020).

References 1. Jyoti B, Faisal K, Rouzbeh A, Vikram G, Roberto O (2015) Modelling of pitting corrosion in marine and offshore steel structures—a technical review. J Loss Prev Process Ind 37:39–62 2. Zou XD, Zhao DP, Sun JC, Wang C, Matsuura H (2018) An integrated study on the evolution of inclusions in EH36 shipbuilding steel with Mg addition: from casting to welding. Metall Mater Trans B 49(2):481–489 3. Kong H, Zhou YH, Lin H, Xia YJ, Li J, Yue Q (2015) Cai ZY (2015) The mechanism of intragranular acicular ferrite nucleation induced by Mg–Al–O inclusions. Adv Mater Sci Eng 6:1–6 4. Pan F, Chen HL, Su Y, Su Y, Hwang W (2017) Inclusions properties at 1673 K and room temperature with Ce addition in SS400 steel. Sci Rep 7(1):2564–2571 5. Wang BX, Liu XH, Wang GD (2018) Inclusion characteristics and acicular ferrite nucleation in Ti-containing weld metals of X80 pipeline steel. Steel Res Int 89(2):2124–2138 6. Jim HH, Shim JH, Cho YW, Lee HC (2015) Formation of intragranular acicular ferrite grains in a Ti-containing low carbon steel. High Temp Mater Process 43(8):1111–1113

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7. Han SK, Chang CH, Lee HG (2005) Evolution of inclusions and resultant microstructural change with Mg addition in Mn/Si/Ti deoxidized steels. Scr Mate 53(11):1253–1258 8. Guo YT, He SP, Chen GJ, Wang Q (2016) Thermodynamics of complex sulfide inclusion formation in Ca-treated Al-killed structural steel. Metall Mater Trans B 47(4):2549–2557 9. Byun JS, Shim JH, Cho YM, Lee DN (2003) Non-metallic inclusion and intragranular nucleation of ferrite in Ti-killed C-Mn steel. Acta Mater 51(6):1593–1606 10. Marie-Aline VE, Guo M, Dekkers R, Burty M, Joris VD (2009) Formation and evolution of Al–Ti oxide inclusions during secondary steel refining. ISIJ Int 49(8):1133–1140 11. Pan F, Zhang J, Chen HL, Su YH, Kuo CL, Su YH, Chen SH, Lin KJ, Hsieh PH, Hwang WS (2016) Effects of rare earth metals on steel microstructures. Materials 9(6):417–436 12. Adabavazeh Z, Hwang WS, Su YH (2017) Effect of adding cerium on microstructure and morphology of Ce-based inclusions formed in low-carbon steel. Sci Rep 7:46503. https://doi. org/10.1038/srep46503 13. Lan J, He JJ, Ding WJ, Wang QD, Zhu YP (2000) Effect of rare earth metals on the microstructure and impact toughness of a cast 0.4C–5Cr–1.2Mo–1.0 V steel. ISIJ Int 40(12):1275–1282 14. Liu HL, Liu CJ, Jiang MF (2012) Effect of rare earths on impact toughness of a low-carbon steel. Mater Des 33(1):306–312 15. Chen X, Li YX (2007) Fracture toughness improvement of austempered high silicon steel by titanium, vanadium and rare earth elements modification. Mater Sci Eng, A 444(1–2):298–305 16. Chen L, Ma XC, Wang LM, Ye XN (2011) Effect of rare earth element yttrium addition on microstructures and properties of a 21Cr-11Ni austenitic heat-resistant stainless steel. Mater Des 32(4):2206–2212 17. Xi XJ, Lai CB, Li JS, Wang ZG, Sun LF, Chen YJ (2017) Effect of Y-base rare earth on the microstructure and impact toughness of E36 steel plate. Chin J Eng 39(2):244–250 (In Chinese) 18. Liu CJ, Huang YH, Jiang MF (2011) Effects and mechanisms of RE on impact toughness and fracture toughness of clean heavy rail steel. J Iron Steel Res Int 18(3):52–58

Part VII

Additive Manufacturing of Metals: Microstructural Evolution and Phase Transformations

Influence of Nitrogen on Microstructure, Mechanical Properties and Martensitic Phase Transformation of Co–26Cr–5Mo–5W Alloys by Selective Laser Melting Bo Wang, Xinglong An, Fei Liu, Min Song, Song Ni and Shaojun Liu Abstract The relationship between the microstructure, mechanical properties and martensitic phase transformation of N-containing Co–26Cr–5Mo–5W alloys by selective laser melting (SLM) is studied. The high-resolution transmission electron microscope and X-ray diffraction observations show that two phases (ε and γ phase) co-exist in N-free alloys. In contrast, a significant decrease of ε phase is observed in N-containing alloys. It is believed that the stabilization of γ phase in N-containing alloys results from the lattice distortion and Si-rich fine-distributed precipitates that block the motion of dislocations. Both the ultimate tensile strength and 0.2% proof strength can be significantly improved by nitrogen addition. Interestingly, the elongation slightly increases as well. They are ~1385 MPa, ~1140 MPa, and ~18.4% for 0.08 N-containing Co–26Cr–5Mo–5W alloys, respectively. It is clear that nitrogen addition during the SLM processing could be a promising strategy to fabricate Co–Cr–Mo–W alloys with an excellent combination of strength and ductility by suppressing face-centered cubic (fcc) → hexagonal close-packed (hcp) martensitic phase transformation. Keywords Co–Cr–Mo–W alloys · Selective laser melting · Nitrogen doping · Martensitic phase transformation

B. Wang · X. An · F. Liu · M. Song · S. Ni · S. Liu (B) Powder Metallurgy Research Institute, Central South University, Changsha 410083, China e-mail: [email protected] S. Liu Central South University, Shenzhen Research Institute, Shenzhen 510085, China © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_40

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Introduction Co–Cr-based alloys have been widely used in dental restoration and orthopedic implants, which include total hip and knee replacements, due to their excellent mechanical properties, corrosion resistance, wear resistance, and biocompatibility [1–3]. It is extremely important to maintain high strength and good ductility for their applications in total hip and knee replacements in orthopedic implant, dental crowns and replaceable dentures [4, 5]. It has been shown that grain sizes, precipitates, volume fraction of γ and ε phase, dislocations, and stacking faults (SFs) significantly influence the mechanical properties of Co–Cr–Mo alloys [6–8]. Previous studies showed that nickel can significantly enhance the ductility of ascast Co–Cr alloys due to the inhibition of γ-fcc phase to ε-hcp martensitic phase transformation [9–11]. Unfortunately, Ni can cause allergies and cancer in living organisms. Further investigation revealed that doping of nitrogen into alloys can improve the ductility of biomedical Co–Cr alloys as well. A small amount of nitrogen addition can induce nano-scaled Cr2 N precipitates or short-range ordering (SRO) that inhibit the dislocation slip, thereby suppressing γ → ε martensitic transformation [12–14]. Selective laser melting (SLM) is considered as one of advanced metal additive manufacturing (AM) technologies applicable to the complex geometry components [15, 16]. Owing to the high demand for personalized customization for patients, more attention has been paid on the selective laser melting (SLM) of Co–Cr-based alloys. However, fundamental understanding of the relationship between the processing, microstructure, martensite phase transformation, and mechanical properties in SLMed Co–Cr–Mo alloys still lacks. In the present study, nitrogen-doped Co–26Cr–5Mo–5W-based alloys with significantly enhanced strength and good ductility are successfully fabricated by selective laser melting. The relationships between SLM-induced microstructure, martensitic transformation, and mechanical properties of N-doped Co–26Cr–5Mo–5W alloys were discussed.

Materials and Methods Two kinds of Co–Cr–Mo–W alloy powders (nitrogen-free and nitrogen-doped) were prepared by gas atomization in a nitrogen atmosphere in the State Key Laboratory for Powder Metallurgy at Central South University. Cr2 N powder was used as the nitrogen source [17]. The particle size distribution and average size are ~16–52 μm and ~29.3 μm, respectively. The chemical compositions of the alloy powders evaluated by an inductively coupled plasma optical emission spectrometer (ICP-OES) were summarized in Table 1. Tensile specimens and bulk specimens were prepared by SLM equipment (Farsoon FS271 M) for mechanical properties testing. The mechanical properties of the specimens were tested by double-column bench material testing machine

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Table 1 Chemical compositions of Co–Cr–Mo–W alloys (mass%) Alloy

Co

Mo

W

N

Ni

Mn

C

N-free

Bal. 25.46

Cr

5.04

5.51

0.018

ε according to the SAED

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pattern. It can be seen clearly that the stacking sequence has changed from FCC phase: …ABCABC… to HCP phase: …ABABAB… in HRTEM image. Figure 3c, d shows that no obvious ε phase but some thin SFs and precipitates can be found in N-doped Co–Cr–Mo–W alloys. These results indicate that both the size and amount of ε phases in N-free alloys are larger than in N-doped alloys, which are consistent with previous XRD results as shown in Fig. 2. The SFs can be easily formed in the present study due to the low stacking fault energy of Co–Cr-based alloys, in which the stacking fault energy can even be negative at room temperature [18]. It is well known that SFs can be regarded as local ε-layers pinning in the γ -fcc matrix [13]. It is reported that the formation of thin ε-layer is caused by three different types of Shockley partial dislocations with different shear directions periodically gliding on every second {111} plane to release the local strain [22]. The stacking sequence in [111] a plane of γ phase has been changed from ABCABC… to ABABAB…. Then, some new ε-layers nucleate closely to the preexisting ones and finally grow together. Eventually, the large ε-phase is formed as shown in Fig. 3b. The results indicate that the N doping has a significant influence in suppressing martensitic transformation by suppressing Shockley partial dislocation slip since the dislocation slip is a key step to generate ε phase. Furthermore, short-range ordering (SRO) or nanoscale Cr2 N precipitates can be generated in the γ matrix in N-doped Co–Cr–Mo alloys, which act as obstacles for Shockley partial dislocations gliding and subsequently increase the energy barrier for γ → ε martensitic transformation [12]. In our study, no obvious SRO and Cr2 N were found except for some precipitates. Figure 4 shows the energy dispersive spectrum (EDS) analysis of N-doped Co–Cr–Mo–W alloy. Two types of precipitates were observed, which appear to be white and black, respectively, as seen in Fig. 4. By comparison, both types of precipitates are Co- and Cr-poor and Si-rich. Besides, white precipitates are W-rich, black precipitates are Mo- and W-poor and Mn-rich. It should be mentioned that element volatilization may occur during SLM process. It was reported that the coarse and elongated precipitates are a possible resembling Si-rich inclusions [16]. The pinning role of the precipitates could be another cause for the γ → ε martensitic transformation blocked. However, further investigation is necessary to clarify the martensitic transformation of Co–Cr-based alloys during SLM processing. Especially, the SLMed alloys are of unique characteristics indicated by a high level of residual stress, heterogeneous metastable microstructures, and nonequilibrium elemental compositions or phase distributions.

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Fig. 4 a Bright-field TEM image of N-doped Co–Cr–Mo–W alloys. b The energy dispersive spectrum (EDS) analysis of N-doped Co–Cr–Mo–W alloy

Conclusion Doping nitrogen, which effectively suppresses the fcc γ phase to hcp ε phase martensitic transformation, significantly enhances the ultimate tensile strength and 0.2% proof strength of Co–Cr–Mo–W based biomedical alloys by SLM while the fracture elongation can be well maintained. The ultimate tensile strength, 0.2% proof strength, and elongation are ~1385, ~1140 MPa, and ~18.4% for 0.08 N-containing Co–26Cr–5Mo–5W alloys, respectively. Combining with SLM, doping of nitrogen into Co–Cr–Mo–W alloys provides an effective strategy to fabricate biomedical Co–Cr–Mo–W-based alloys with extremely high strength and good ductility.

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Acknowledgements The use of facilities in the State Key Laboratory for Powder Metallurgy and the Institute for Materials Microstructure at Central South University is acknowledged.

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The Morphology, Crystallography, and Chemistry of Phases in Wire-Arc Additively Manufactured Nickel Aluminum Bronze Chalasani Dharmendra, Amir Hadadzadeh, Babak Shalchi Amirkhiz and Mohsen Mohammadi

Abstract A new Wire-Arc Additive Manufacturing (WAAM) technique is used to produce Nickel Aluminum Bronze (NAB) components for marine applications in view to mitigate the problems that typically arise in a cast microstructure. In cast condition, the alloy typically exhibits microstructure that consists of an FCC Cu-rich solid solution (or α-phase), some retained β-phase, and several intermetallic phases collectively referred to as κ-phase. This study aims to characterize the crystal structures of the various κ-phases or precipitates, their distribution, morphology, orientation relationships with the α-matrix, and their chemical compositions in WAAM-NAB alloy using electron microscopy. The precipitation of κ-phase differs in morphology and chemical composition to those present in a cast NAB. In addition, some uniaxial tensile coupons were machined out of the WAAM-NAB samples, where tensile mechanical properties are superior to those of cast NAB. The effects of microstructural differences in both alloys on the mechanical properties are correlated. Keywords Nickel aluminum bronze (NAB) · Additive manufacturing (wire-arc) · Electron microscopy · Microstructure · Mechanical properties

Introduction Owing to their good combination of strength, fracture toughness, corrosion resistance, friction coefficients, and non-sparking behavior, Nickel Aluminum Bronze (NAB) alloys are widely used for marine, shipbuilding, and offshore oil and gas components [1]. The nominal composition of the NAB alloy studied here is 8.5–11 wt% C. Dharmendra (B) · A. Hadadzadeh · M. Mohammadi Marine Additive Manufacturing Centre of Excellence (MAMCE), University of New Brunswick, Fredericton, NB, Canada e-mail: [email protected] A. Hadadzadeh · B. S. Amirkhiz CanmetMATERIALS, Natural Resources Canada, Hamilton, ON, Canada © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_41

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Al, 4–5 wt% Ni, 2–5 wt% Fe, 0.8–1.5 wt% Mn, and the balance is Cu. The addition of nickel and iron to the binary Cu–Al alloys has been found to increase the mechanical properties of the NAB through the precipitation of complex κ-phases in both the α and β phase fields [2]. Both Ni and Fe extend the terminal FCC α-phase field and suppress the formation of γ-phase that corrodes preferentially, especially when the Al content is more than 9.5 wt% in the Cu–Al system [3]. NAB alloys are traditionally produced via casting, where the Mn increases the fluidity of the molten alloy. The microstructure under normal casting conditions (coupled with slow cooling rates for thick sections) typically consists of coarse Cu-rich α-phase (FCC), retained β phase during cooling from high temperatures, and various Ni–Fe–Al intermetallic phases (precipitates) denoted as κ-phases [4]. Porosity is an inevitable issue during casting, which reduces the physical and mechanical properties. To improve the microstructure and mechanical properties, several other methods such as heat treatment, friction welding, and friction stir processing (FSP) are reported in the literature. In comparison to the casting technology, the additive manufacturing (AM) technology offers a great advantage by saving cost and time, especially to produce components in complex shapes, as it is a layerby-layer manufacturing technique. Wire-Arc Additive Manufacturing (WAAM) is unique as it combines the use of wire as depositing material and electric arc as a fusion source to fabricate the components to overcome the common defects that arise during casting. At present, there is very limited work [5] reported on the feasibility of producing NAB components using the WAAM processes, but there is no information on the microstructural development of additively manufactured NAB alloy. The present study is the first of its kind and the object is to investigate the morphology and chemistry of complex phases formed using optical and electron microscopy techniques as well as comparing with cast microstructure and correlating with room temperature mechanical properties.

Materials and Methods Berlin-based GEFERTEC has invented a new-patented technology, named as 3DMP process that uses a wire feedstock. This system uses an electric arc welding, in which a wire that passes through the feedstock melts and deposits in successive layers. CAD model was used and the CAM software directed the CNC arc welding head to deposit layers. A wire that has a designation of C95800 with 8.5–9.5 wt% Al, 4–5 wt% Ni, 3.5–4.5 wt% Fe, 0.8–1.5 wt% Mn, and Cu (balance) was used in the WAAM process. Figure 1a shows the NAB alloy that was printed in vertical direction (15 cm in length) and in cuboidal form (two other sides are in 2.5 cm) on a Stainless steel substrate. Cast NAB alloy from a pump impeller is shown in Fig. 1b. The microstructure of the as-cast and WAAM-NAB samples was analyzed using optical microscopy, scanning electron microscopy (JEOL JSM6400 SEM), and transmission electron microscopy (FEI Tecnai Osiris TEM). EDS X-ray detection system in the scanning transmission electron microscopy (STEM) mode was used to

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Fig. 1 a Wire-arc additively manufactured (WAAM) and b cast nickel aluminum bronze alloy

analyze the precipitates. Samples for TEM were prepared using Gatan 691 PIPS ion milling equipment. The OM and SEM samples were polished for a mirror-like surface by standard mechanical polishing procedures. Then, the samples were etched with Klemm’s reagent (mixture of sodium thiosulfate solution and potassium metabisulfite). Room temperature tension tests on both WAAM and cast NAB were performed on an INSTRON 1332 model at a strain rate of 0.001 s−1 .

Results and Discussion Microstructure of the WAAM-NAB Alloy and Comparison with Cast NAB The microstructure of the WAAM-NAB is shown in Fig. 2a. The NAB part produced using the WAAM process has no porosity and reached to full density as can be seen by the layer bands in Fig. 1a. The microstructure consists of a bright α-phase (Curich), the dark-etched martensitic β phase (retained during solidification or phase transformation), and intermetallic particles referred as κ-phases. For the comparison purpose, optical micrograph of the cast NAB alloy is also shown in Fig. 2b. The large white areas correspond to the α-phase. Small areas of β phase and other intermetallic phases were etched in black contrast. Scanning electron micrographs (SEM) of cast and the WAAM-NAB alloy are presented in Fig. 3a and 3b, respectively, which clearly reveals various particles in the microstructure.

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Fig. 2 Optical microstructures of a WAAM (as-printed) and b cast NAB alloy

κ-Phases Dendritic or rosette-shaped particles that are located mainly in the α-grain centers are κI precipitates, as marked in Fig. 3a for cast NAB. These particles were identified as iron-rich precipitates and each particle is composed of a number of different structures including disordered iron-rich solid solution (BCC), Fe3 Al (DO3 ), and FeAl (B2) [6]. The κII precipitates are similar in the shape but are smaller than the κI precipitates.

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Fig. 3 SEM microstructures of a cast NAB and b WAAM (as-printed) alloy showing constituent phases

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These globular precipitates are mainly distributed at the grain boundaries with a size of about 2–3 μm and were identified as Fe3 Al based. The κIII phase is a product of eutectoid transformation, which appears as lamellar (as shown in Fig. 3a) or globular (degenerate lamellar structure). These precipitates were found to be rich in nickel and aluminum [7]. A dispersion of fine κIV precipitates can be seen (Fig. 3a) within the α-phase. In addition, there were some precipitate free zones at the periphery of the α-grains. Figure 3b is the SEM image showing the microstructure of additively manufactured NAB alloy. Compared with the cast NAB, instead of κI and κII phases, strips of continuous and coarse κIII intermetallic compound (NiAl) are formed at the grain boundary of α-phase as well as at the periphery of the lamellar eutectoid regions. Differentiating the β phase from the NiAl lamellar’s in the SEM image is hard due to the negligible contrast difference. The reduction in volume fraction of Fe–Al-based κ phase in the WAAM-NAB is due to the fast cooling rates that suppress the eutectoid reaction β → α + κ. In order to better understand the morphologies of the precipitates, TEM analysis was performed. Figure 4a shows the HAADF-STEM image of the WAAM-NAB microstructure along with its superimposed energy-dispersive spectroscopy (EDS) compositional map that is shown in Fig. 4b. Figure 4c–g presents the EDS elemental maps for Cu, Al, Ni, Fe, and Mn, respectively. Figure 4h shows five locations of the phases from which EDS compositional analysis was obtained and the results in wt% are listed in Table 1. Most of the Ni–Al-rich regions are in lamellar and are likely the κIII phase along with the β phase (martensitic) that is retained β during solidification from the β region. A bright spherical region at the center is Fe–Mn rich at the core (point 1 in Table 1) surrounded by Ni–Al-rich precipitate in globular form at the periphery. Figure 4f illustrates spheroidal morphologies of fine precipitates that are believed to be Fe-rich κIV . The intensity or the presence of Fe seems to differentiate the κIII (NiAl, Fe) with the β (right-side top region). The temperature at which κIV formation begins during solidification is influenced by the temperature of α formation and by the Fe/Ni ratio. The higher the temperature at which α forms, the higher the iron content is likely to be, and consequently, the precipitation of κIV . The regions that form first from β-phase are centers of the grains of α-phase, which contain highest amount of iron in solid solution. Centers of α-grains are the first regions in which the limit of solid solubility will be reached during cooling and this variation in iron content across the α-grains is reflected in the form of κIV precipitate free zones, as shown in Fig. 4a.

The Martensitic Phase Martensite phase (β ) is actually ‘retained β’ phase, which is considered as undesirable because of its susceptibility to preferential corrosion [8]. But it is difficult to avoid this phase formation even under slow cooling rates that are typically involved in casting. In the cast NAB, it was reported that this phase has complex structure

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Fig. 4 High-angle annular dark-field (HAADF)-STEM image of WAAM-NAB that shows a microstructure b combined EDS map and the corresponding elemental maps for c Cu, d Al, e Ni, f Fe, g Mn and h five points from which the EDS spectrums were collected

with high density of NiAl-based precipitates. This phase is formed due to the segregation of Al, and high content of Al stabilizes the β phase to lower temperatures at which martensitic transformation starts. Before the formation of κIII particles, the Ni and Al contents in the β phase (α + β region in the phase diagram at solidus temperature) are at their peak. During the decomposition of β into α + κIII , Ni forms compounds with Al (NiAl) instead of completely going into solid solution, while the remaining Al stabilize the β phase. This can be avoided only if cooling time is long enough to allow complete eutectoidal decomposition of β. As additive manufacturing techniques (layer-by-layer deposition) typically employ high cooling rates, the

450 Table 1 Composition of the various phases or precipitates obtained by EDS from Fig. 1h

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Element

Composition (Weight %) Point 2

Point 3

Point 4

Point 5

Cu

Point 1 2.21

63.66

89.18

56.71

85.09

Al

0.54

–

6.61

14.94

6.35

Ni

0.63

16.21

1.93

18.48

3.98

Fe

73.75

5.64

1.04

8.05

3.19

Mn

8.51

1.71

1.21

1.8

1.37

Cr

10.36

–

–

–

–

V

0.71

–

–

–

–

Fig. 5 TEM micrograph of the WAAM-NAB alloy showing the martensite phase

presence of retained β (β ) is expected in the final microstructure of WAAM-NAB alloy. Figure 5 shows the TEM micrograph that reveals the presence of martensitic phase in the WAAM-NAB alloy.

Tensile Testing The room temperature tensile properties of the WAAM-NAB alloy along with cast NAB are shown in Fig. 6. Compared to the cast alloy, additively manufactured NAB alloy exhibits similar yield strength, relatively higher tensile strength, and more importantly superior ductility. The variations in mechanical properties are directly connected to the cooling rate, chemical composition, and the resultant microstructure of the alloys. Among the phases in cast NAB alloy, the α-phase matrix is a soft Curich solid solution that has high ductility, whereas the martensitic phase β is stronger and harder [9]. The precipitates including globular or rosette-shaped κII and lamellar

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Fig. 6 Engineering stress–strain curves of the WAAM (3D printed) and cast NAB samples

κIII are hard phases and less ductile. The improvement in strength and ductility for WAAM-NAB alloy is attributed to the change in ratio of phases (i.e. no κI and κII ) and their morphology, reduced size of α matrix. The mechanical properties of multiphase materials are a result of comprehensive effect of number of factors such as the presence of various solutes and their concentration, shape, size, and distribution in the matrix. More in-depth study is required to understand the strengthening and/or deformation mechanisms in the WAAM-NAB compared to the cast alloy. The general observation on the mechanical properties of AM alloys is that the tensile strengths show relatively higher value in the longitudinal and transverse print directions than in the normal direction [10]. The WAAM-NAB alloy in the present case can be referred as the alloy printed in the normal direction. The strength of this alloy may be even higher than the cast alloy in case if it is printed in longitudinal direction.

Fractography Figure 7a, b shows the fracture morphologies of the WAAM and cast NAB alloys after tensile tests. The SEM micrograph of fracture surface (Fig. 7a) shows more number of conical equiaxed dimples suggesting a good ductile fracture nature of the WAAM-NAB alloy. As shown in Fig. 7b, the fracture morphology of cast NAB is similar to that of the WAAM-NAB, except that the dimples are less in quantity.

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Fig. 7 Tensile fracture morphologies of the a WAAM and b cast NAB samples

Conclusion The intent of this work has been to explore the feasibility of adopting wire-arc additive manufacturing to produce nickel aluminum bronze alloy, and to study the morphology and the distribution of the intermetallic phases using optical, electron, and transmission electron microscopy techniques. The findings are drawn as follows: 1. A nearly full dense sample of NAB without any apparent porosity even at microscopic level was built through WAAM process in cuboidal shape of 15 cm height and 2.5 cm width. 2. The as-cast microstructure consists of copper-rich α phase, retained β, and various κ-phases, based on Fe3 Al or NiAl. The Fe-rich κI and κII phases were suppressed in the WAAM-NAB alloy. 3. BCC β phase that did not undergo complete decomposition on cooling due to fast cooling rates in WAAM process transformed into a complex martensitic structure containing high density of NiAl precipitates. 4. The ductile fracture mode was determined to be the main failure mode of both NAB alloys. WAAM-NAB exhibited better ductility and slightly higher strength than the cast NAB. Acknowledgements The authors would like to acknowledge Alexander Riemann of Gefertec GmbH, Berlin, Germany for printing the nickel aluminum bronze via WAAM technique at their facility. The authors would like to thank Catherine Bibby of CanmetMATERIALS, Hamilton, Canada for TEM samples preparation.

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References 1. Duma JA (1975) Nav Eng J 87:45–64 2. Brezina P (1982) Int Met Rev 27:77–120 3. Hansen M, Anderenko K (1958) Constitution of Binary Alloys, 2nd edn. NcGraw-Hill, New York, NY, pp 84–89 4. Culpan EA, Rose G (1978) J Mater Sci 13:1647–1657 5. Ding D, Pan Z, Duin SV, Li H, Shen C (2016) Materials 9:652–663 6. Hasan F, Jahanafrooz A, Lorimer GW, Ridley N (1982) Metall Trans A 13A:1337–1345 7. Jahanafrooz A, Hasan F, Lorimer GW, Ridley N (1983) Metall Trans A 14A:1951–1956 8. Weill-Couly P, Arnaud D (1973) Fonderie 28:123–135 9. Wu Z, Cheng YF, Liu L, Lv WJ, Hu WB (2015) Corros Sci 98:260–270 10. Asgari H, Baxter C, Hosseinkhani K, Mohammadi M (2017) Mater Sci Eng A 707:148–158

Microstructure Evolution in Direct Metal Laser Sintered Corrax Maraging Stainless Steel Amir Hadadzadeh, Babak Shalchi Amirkhiz, Jian Li and Mohsen Mohammadi

Abstract Martensitic hardenable (maraging) stainless steels are of interest due to their combination of high strength and ductility along with superior corrosion and stress corrosion cracking properties. Corrax maraging steel has been recently atomized in powder form for laser-based sintering applications. In the current study, Corrax with a nominal composition of 11–13% Cr, 8.4–10% Ni, 1.1–1.7% Mo, 1.2–2% Al and 0.05% C (in wt%) was additively manufactured through direct metal laser sintering (DMLS) process. This additive manufacturing technique results in ultrafine microstructures with unique micron and submicron characteristics. The microstructure of the as-built DMLS-Corrax was investigated using SEM, EBSD, and TEM. The result of the current study is the preliminary step to develop additively manufactured high strength Corrax stainless steels for various applications. Keywords Additive manufacturing · Direct metal laser sintering (DMLS) · EBSD · TEM · Maraging steels

Introduction There is an increasing demand in commercial and military shipbuilding industries for simultaneous improvement of both strength and ductility of the structural materials to enhance the performance of the vessels [1]. Meanwhile, the corrosion properties of the materials are of importance for the marine industries since corrosion is a serious problem in such applications [2]. Amongst the available structural materials, maraging steels exhibit the desirable properties (i.e. ultrahigh strength and good A. Hadadzadeh (B) · M. Mohammadi Marine Additive Manufacturing Centre of Excellence (MAMCE), University of New Brunswick, Fredericton, NB E3B 5A1, Canada e-mail: [email protected] A. Hadadzadeh · B. S. Amirkhiz · J. Li CanmetMATERIALS, Natural Resources Canada, 183 Longwood Road South, Hamilton, ON L8P 0A5, Canada © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_42

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ductility) for marine applications due to possession of relatively ductile martensitic matrix (with very low carbon content), which is strengthened by nanoscale intermetallic precipitates [3]. Conventional maraging steels contain a high amount of Ni (usually around 18 wt% [4]) and other elements such as Co (8–13 wt%) and Mo (3–5 wt%) [5]. High Ni content promotes the formation of martensite during decomposition of austenite; however, Ni-enriched martensite is not thermally stable and transform to austenite at high temperature [5]. To overcome this challenge, a lower amount of Ni (8–10 wt%) is added to the chemical composition, which makes NiAl as the strengthening intermetallic precipitate [6] (instead of Ni3 Ti in the traditional maraging steels). Moreover, the addition of high Cr content (up to 17 wt%) improves the corrosion resistance [7]. PH13–8 Mo, a precipitation-hardening (PH) martensitic stainless steel, is one of the most famous materials in low Ni content maraging steels family [8]. These steels obtain their maximum strength through solution annealing followed by aging heat treatment. Due to disruption of additive manufacturing (AM) techniques as revolutionary manufacturing routes [9], it is imperative to develop the required knowledge on how different materials, including maraging steels, can be used for fabrication of AM parts. Unlike the high carbon content steels that face cracking and pores formation during additive manufacturing [10], low carbon content ones are potentially compatible for AM processes [11]. Recently, a new class of precipitation-hardening stainless steel (called CX) in the family of PH13–8 Mo has been developed by EOS GmbH [12] for fabrication of AM parts through direct metal laser sintering (DMLS). This material is almost the same as Uddeholm-Corrax® [10]. Asgari and Mohammadi [10] showed that DMLS-Corrax samples could achieve yield and ultimate tensile strengths of 1036 MPa and 1113 MPa, respectively, in the as-built condition. Therefore, this alloy exhibits high strength, even in the as-built condition without any further heat treatment. In order to understand how the microstructure of DMLS-Corrax develops in the as-built condition, cuboid samples of this material were fabricated through DMLS process. The as build microstructure was then analyzed using scanning electron microscopy (SEM), electron backscatter diffraction (EBSD) and transmission electron microscopy (TEM).

Experimental Procedure Cuboid samples with dimensions of 15 mm × 15 mm × 15 mm (shown in Fig. 1) were fabricated through DMLS process using Corrax powder with the nominal chemical composition of Fe–12 wt%Cr–9.2 wt%Ni–1.4 wt%Mo–1.6 wt%Al–0.4 wt%Mn–0.4 wt%Si–0.05 wt%C. The powder particles were mainly spherical or near spherical in morphology with an average particle size of 37.5 ± 16 µm. The samples were made using an EOS M290 machine equipped with a 400 W Yb-fibre laser (maximum capacity) and employing the process parameters developed by EOS GmbH to achieve the least porosity as reported elsewhere [10]. The building plate was preheated and

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Fig. 1 Schematic of cuboid sample along with the orientation of samples for microstructural analysis. Note: “z” is the building direction

maintained at 80 °C to minimize the residual stresses and a stripe scanning strategy where the laser beam rotated 67° between the successive layers was employed. The microstructure of the as-built sample was analyzed using EBSD over the top and side of the sample (as shown in Fig. 1) using a field emission gun scanning electron microscope (FEG-SEM FEI Nova NanoSEM-650) equipped with a Hikari EBSD system. The scans were conducted at three magnifications of × 300, × 2000 and × 5000 over an area of 250 µm × 650 µm, 100 µm × 100 µm and 40 µm × 40 µm, respectively. The scan step size was set as 0.3 µm, 0.07 µm and 0.05 µm for the three magnifications, respectively. SEM studies were conducted using the same system on a lightly etched sample. The sample was etched for 5 s with Kalling’s reagent (5 g CuCl2 -100 ml HCl-100 ml ethanol). Details of microstructure were studied through TEM using an FEI Tecnai Osiris TEM equipped with a 200 keV X-FEG gun.

Results and Discussion Figure 2 shows the SEM microstructure of DMLS-Corrax from the top view. The microstructure consists of packets of fine lath structure, distributed in random directions. The lath structure is similar to previously reported lath-like ferrite [13], lath bainite [14] and lath martensite [5, 15]. Since the sample was not heat treated, it is doubtful that martensite developed in the microstructure. On the other hand, the fast cooling rates during DMLS process in the range of 103 –108 K/s [9] can facilitate the prior austenite to martensite transformation. However, it is not possible to recognize the phases from the SEM microstructure.

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Fig. 2 SEM microstructure of DMLS-Corrax from top view

Figure 3 shows the typical EBSD maps of the DMLS-Corrax sample from the top and side views, at a magnification of × 300. From the top view, the laser scan tracks featured by elongated grains are visible. The elongated grains were resulted from the epitaxial growth at the fusion boundaries [16]. The microstructure of DMLS-Corrax is featured by fine grains from the side view (along the building direction) with an average size of 5.2 ± 1.5 µm. The grains are very uniform in size so that it is not possible to recognize the melt pool boundaries. Moreover, the grains were analyzed considering the grain shape aspect ratio (φ  L 2 /L 1 , where L 1 and L 2 are the grain major and minor axes, respectively). L 1 and L 2 were evaluated by fitting an ellipse to each grain and the columnar and equiaxed grains were recognized with φ ≤ 0.33 and φ > 0.33, respectively [17]. DMLS-Corrax sample consisted of 72% equiaxed and 28% columnar grains. Details of the grain structure along the building direction are shown at a magnification of × 2000 in Fig. 4. The lath structure is observed clearly in the EBSD map. The high angle grain boundaries (HAGBs) with a misorientation of > 5° were marked as black lines and superimposed on the same map in Fig. 4. All the lath boundaries coincided with the HAGBs. However, it is not possible to recognize the prior austenite grain boundaries or packet boundaries from the EBSD map. To analyze the amount of retained austenite in the microstructure of DMLSCorrax, phase analysis was conducted at a high magnification of × 5000, and a typical

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Fig. 3 EBSD unique color grain maps of DMLS-Corrax from a top view and b side view

result from the top view is shown in Fig. 5. The HAGBs map was superimposed on the phase map. The matrix phase was marked as ferrite in the map, since ferrite, bainite and martensite are detected as a single phase through the EBSD diffraction patterns. As seen in the EBSD phase map, around 2% of the microstructure is consisted of retained austenite due to incomplete transformation of austenite. Almost all the retained austenite grains are observed along the HAGBs. STEM bright field (STEM-BF) microstructure of DMLS-Corrax from top view is observed in Fig. 6. The lath structure without any noticeable precipitation along with the retained austenite is observed in the microstructure. Moreover, the laths are featured by high density of dislocations which has been reported for both lath bainite [14] and lath martensite [5]. It appears that all the laths shown in Fig. 6 belong to a single packet, since all the laths are parallel with almost the same thickness. Although the laths observed in the current study are similar to those observed for bainite and martensite, but it is not possible to recognize them from the current observations. More in-depth analyses are required in future studies to clearly identify the lath phase.

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Fig. 4 EBSD unique color grain map of DMLS-Corrax from side view

Fig. 5 EBSD phase map of DMLS-Corrax from top view, superimposed with the HAGBs map

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Fig. 6 STEM-BF microstructure of DMLS-Corrax from top view. RA stands for retained austenite

Conclusions In the current study, cuboid samples of a low Ni content maraging steel (Corrax) were additively manufactured through DMLS process. The microstructure of the as-built sample was analyzed over the top and side of the sample using SEM, EBSD and TEM techniques. It was observed that the as-built microstructure was consisted of fine grains with an average size of 5.2 ± 1.5 µm. The majority of the grains were equiaxed, decorated with the retained austenite. High-resolution EBSD studies revealed the lath structure of the matrix, which also was confirmed through TEM analysis. The laths were featured by high density of dislocations; however, it was not possible to recognize them (ferrite, bainite or martensite) in the current study. Therefore, more in-depth analyses will be conducted in future to clarify the lath phase. Acknowledgements The authors would like to acknowledge Natural Sciences and Engineering Research Council of Canada (NSERC) project number RGPIN-2016–04221 and New Brunswick Innovation Foundation project number (NBIF)-RIF2017–071 for the financial support of this work. The authors would also like to acknowledge Dr. Mark Kozdras at CanmetMATERIALS for facilitating the research and Pei Liu for TEM sample preparations.

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References 1. Cao X, Wanjara P, Huang J, Munro C, Nolting A (2011) Hybrid fiber laser—Arc welding of thick section high strength low alloy steel. Mater Des 32:3399–3413 2. Guedes S C, Garbatov Y, Zayed A, Wang G (2009) Influence of environmental factors on corrosion of ship structures in marine atmosphere. Corros Sci 51:2014–2026 3. Simm TH, Sun L, Galvin DR, Gilbert EP, Alba Venero D, Li Y, Martin TL, Bagot PAJ, Moody MP, Hill P, Bhadeshia MKDH, Birosca S, Rawson MJ, Perkins KM (2017) A sans and apt study of precipitate evolution and strengthening in a maraging steel. Mater Sci Eng A 702:414–424 4. Tewari R, Mazumder S, Batra IS, Dey GK, Banerjee S (2000) Precipitation in 18 wt% Ni maraging steel of grade 350. Acta Mater 48:1187–1200 5. Sun L, Simm TH, Martin TL, McAdam S, Galvin DR, Perkins KM, Bagot PAJ, Moody MP, Ooi SW, Hill P, Rawson MJ, Bhadeshia HKDH (2018) A novel ultra-high strength maraging steel with balanced ductility and creep resistance achieved by nanoscale b-NiAl and Laves phase precipitates. Acta Mater 149:285–301 6. Guo Z, Sha W, Vaumousse D (2003) Microstructural evolution in a PH13–8 stainless steel after ageing. Acta Mater 51:101–116 7. Seetharaman V, Sundararaman M, Krishnan R (1981) Precipitation hardening in a PH 13–8 Mo stainless steel. Mater Sci Eng 47:1–11 8. Leitner H, Schnitzer R, Schober M, Zinner S (2011) Precipitate modification in PH13–8 Mo type maraging steel. Acta Mater 59:5012–5022 9. DebRoy T, Wei HL, Zuback JS, Mukherjee T, Elmer JW, Milewski JO, Beese AM, WilsonHeid A, Ded A, Zhang W (2018) Additive manufacturing of metallic components–—Process, structure and properties. Prog Mater Sci 92:112–224 10. Asgari H, Mohammadi M (2018) Microstructure and mechanical properties of stainless steel CX manufactured by direct metal laser sintering. Mater Sci Eng A 709:82–89 11. Liu L, Ding Q, Zhong Y, Zou J, Wu J, Chiu YL, Li J, Zhang Z, Yu Q, Shen Z (2018) Dislocation network in additive manufactured steel breaks strength–ductility trade-off. Mater Today 21:354–361 12. EOS (2017) GmbH—Electro Optical Systems, Material Data Sheet: EOS Stainless Steel CX, München, www.eos.info 13. Liang J, Zhao Z, Tang D, Ye N, Yang S, Liu W (2018) Improved microstructural homogeneity and mechanical property of medium manganese steel with Mn segregation banding by alternating lath matrix. Mater Sci Eng A 711:175–181 14. He SH, He BB, Zhu KY, Huang MX (2017) On the correlation among dislocation density, lath thickness and yield stress of bainite. Acta Mater 135:382–389 15. Kitahara H, Ueji R, Tsuji N, Minamino Y (2006) Crystallographic features of lath martensite in low-carbon steel. Acta Mater 54:1279–1288 16. Kou S, (2003) Welding Metallurgy. 2nd ed. Wiley 17. Hadadzadeh A, Amirkhiz BS, Li J, Mohammadi M (2018) Columnar to equiaxed transition during direct metal laser sintering of AlSi10 Mg alloy: effect of building direction. Addit Manuf 23:121–131

The Microtexture and Tensile Properties of Continuous-Wave and Quasi-Continuous-Wave Laser Powder-Deposited Inconel 718 Zhaoyang Liu, Qiang Zhu and Lijun Song

Abstract Inconel 718 superalloy has been widely used in the aerospace field. In this study, the effect of linear heat input and laser mode on the microtexture and tensile properties of laser powder-deposited (LPD) Inconel 718 was studied. Continuouswave (CW) and quasi-continuous-wave (QCW) laser modes were conducted in the experiments. The results showed that under low linear heat input, the microtexture morphologies with CW LPD and QCW LPD were long directional columnar grains and chaotic grains, respectively. The grain size of QCW LPD is finer than that of CW LPD. With the increase of linear heat input, the directivity of microtexture morphology with CW LPD becomes unobvious, while the microtexture morphology of QCW LPD changes little with enlarged grain size. Due to the finer grain size, the tensile strength at 600 °C of sample with low linear heat input (60 J/mm) and QCW laser mode is 17.7% higher than that with high linear heat input (120 J/mm) and CW laser mode. Keywords Microtexture · Tensile property · Laser additive manufacturing

Z. Liu (B) College of Innovation and Entrepreneurship, Southern University of Science and Technology, Shenzhen 518055, China e-mail: [email protected] Q. Zhu Department of Mechanical and Energy Engineering, Southern University of Science and Technology, Shenzhen 518055, China Z. Liu · Q. Zhu Shenzhen Key Laboratory for Additive Manufacturing of High-Performance Materials, Southern University of Science and Technology 518055, Shenzhen, China L. Song Hunan Provincial Key Laboratory of Intelligent Laser Manufacturing, Hunan University, Changsha 410082, China © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_43

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Introduction Laser powder deposition (LPD) process, which allows rapid and accurate addition of controlled amounts of material at the required locations with low heat input, is a very promising and effective additive forming technique capable to produce near-netshape dense components of various kinds of material [1–3], and has been successfully and widely used to fabricate high-performance aircraft components with reasonable cost and saving of high-value materials [4–7]. Inconel 718, as a well-known nickel-based superalloy, has been widely used as the material of turbine disks, blades and shafts due to its superior creep and fatigue strength, corrosion resistance at elevated temperatures up to 650 °C [8]. Inconel 718 is strengthened by the D022 ordered body-centered tetragonal phase γ (Ni3 Nb) and ordered L12 intermetallic phase of γ (Ni3(Al, Ti)) [9]. The macro-segregation caused by niobium segregation generally could not be avoided in cast and wrought products. Comparing with the traditional manufacturing methods such as casting and forging, LPD has obvious advantages in refining microstructure and suppressing macro-segregation due to its rapid solidification process [10] and has been used to produce Inconel 718 components directly. While the solidification condition in the molten pool of LPD is influenced by many processing parameters. It is conceivable that the parameters that influence the thermal cycle during LPD might also influence the formation and morphology of these detrimental phases, making it possible to tailor a microstructure with good properties by controlling the parameters. Geometric accuracy and material properties are the two basics, but challenging requirements for the final implementation of laser-metal deposition technology in the aircraft engine industry. The crystallographic microstructure of the material, i.e., the collective nature of the grain orientations, depends on the processing conditions and strong texture often has significant influence on material properties. Many material properties such as Young’s modulus, Poisson’s ratio, tensile strength and ductility are texture specific. The influence of texture on material properties can be as high as 20–50pct [11]. The study and interpretation of texture is fundamentally important toward the ultimate goal of controlling the macroscale behavior of laser-deposited materials. The texture of laser-deposited alloys can be tuned by deposition parameters, such as laser energy, scanning speed, pulse duration, scanning path patterns, and subsequent heat treatment methods. Blackwell [12] studied the tensile properties of laser-deposited Inconel 718 by the laser engineered net shaping system and found the deposit exhibits strongly anisotropic properties. A subsequent hot isotropic pressure treatment can not only significantly reduce the anisotropic ductility of the deposit but also increase the crossbonding strength. Ocelik et al. [13] observed strong fibrous solidification texture in the laser clad tracks of Co-based alloys using orientation imaging microscopy. They discovered that the average grain orientation angle increases substantially with an increasing scanning speed as a result of the changing average local solidification front directions. Moat et al. [14] investigated the variation of melt pool geometry and the resultant texture orientations of laser-deposited Waspaloy with a pulsed diode

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Table 1 Composition of SX superalloy (wt%) Cr

Co

Mo

Fe

Si

Nb

Al

Ti

C

Ni

18.1

0.17

3.1

17

0.08

5.17

0.53

0.89

0.03

Bal.

laser. They found that long cycle periods and short pulse lengths (low duty cycles) are concluded to result in a highly columnar microstructure, while short pulse length reduces the tilt angle of the epitaxial fibrous texture. Moreover, thick-wall manufacturing exhibited less variation in grain size and averaged aspect ratio compared to thin-walled structures that can be explained by lower thermal gradients. Huang et al. [15] studied the microstructures and tensile properties of laser-deposited IN718 through single direction raster scanning and cross direction raster scanning deposition paths, respectively. Similar tensile results were demonstrated while the later path pattern produces higher ductility owing to its more uniform grain size distribution after a full solution heat treatment. Lambarri et al. [16] obtained an ultimate tensile strength of 1380 MPa on the fully aged Inconel 718 samples. The difference between laser-deposited and the forged material properties are attributed to the number of twin boundaries. Investigating the microtexture and resultant mechanical properties in laser-deposited material can provide significant guidance for better control of the resulting material performance, while the coupling effect of linear heat input and laser mode on the microstructure and tensile property has not been investigated in detail. In this study, the microtexture and tensile properties of Inconel 718 manufactured by continuous-wave (CW) and quasi-continuous-wave (QCW) laser powder deposition process are investigated. The effects of linear heat input on the microtexture and tensile properties of CW LPD and QCW LPD Inconel 718 are studied.

Experimental Procedure A commercially available Inconel 718 powder manufactured by plasma rotating electrode process was used in this study. The chemical compositions of the Inconel 718 powders are shown in Table 1. The particle size of Inconel 718 powder was 65–110 μm. The powder particles are mostly spherical in shape, as shown in Fig. 1. Casting Inconel 718 superalloy plates with size of 100 mm × 100 mm × 20 mm in X, Y and Z directions, respectively, were used as the substrates. The surface of the substrate was grounded with 300-grit SiC paper and cleaned by acetone and alcohol before laser experiments. Figure 2 presents the schematic illustration of LPD experiments. A Ytterbium fiber laser with maximum 2 KW power was used as the heat source for all experiments. The laser beam has a Gaussian model intensity profile. The powder was fed into the molten pool through the annular conical channel of coaxial nozzle. The LPD process is protected by argon shielding gas (purity of 99.99%) fed through the coaxial nozzle

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Fig. 1 Morphology of Inconel 718 alloy powder

with a flow rate of 10 L/min. The laser beam has a near Gaussian model intensity profile and generates a 2 mm circular spot on the surface of samples with a−6 mm defocusing length. The carrier gas of powder is 6 L/min. Thin-wall samples (width 50 mm, height 50 mm, thickness 2 mm) were fabricated for making the tensile test specimens. The thin-wall structures are made by uni-direction scanning paths. Test specimens are designed to be cut out along the vertical direction of the laserdeposited thin-walled samples. Table 2 shows the processing parameters used in the experiments. The per layer height was constantly set as 0.23 mm. For QCW laser mode, the laser power was modulated by a square wave with a duty cycle of 50% and a modulation frequency of 50 Hz with on-power of 480 W and off-power of 0 W. The linear heat input is defined as P/V , where P is the laser power and V is the scanning speed. The samples were subjected to a standard heat treatment method of direct aging which is scheduled follows: heat to 991 K (718 °C), hold for 8 h, furnace cool to 894 K (621 °C), hold for 10 h, and air cool. Test specimens were then machined by wire electro-discharge machining. The samples for microstructure analysis were sectioned, mounted and polished using standard metallographic techniques for nickel-based superalloys, and etched by a solution containing CuCl2 (5 g), HCl (100 mL) and CH3 CH2 OH (100 mL). The microstructure was investigated by an optical microscope and. All the tensile tests were conducted at 873 K (600 °C) in tensile testing machine.

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Fig. 2 Schematic illustration of LPD experiments Table 2 Main material properties for substrate and powder Samples

Laser power (W)

Scanning speed (mm/s)

Power feeding rate (g/min)

1

480

8

9

2

480

8

3

480

4

4

480

4

Linear heat input (J/mm)

Laser mode

60

CW

9

60

QCW

9

120

CW

9

120

QCW

Results and Discussion Figure 3 shows the microstructure in the laser-deposited Inconel 718. As shown in Fig. 3a, under the low linear heat input (60 J/mm), long and coarse columnar dendrites form in the CW LPD sample. The crystallographic orientation of coarse columnar dendrites uniformly tends to the laser scanning direction. This is because during CW LPD, heat is mainly dissipated through the substrate or below layers. The cooling rate at the solidification interface is so large that columnar dendrites grow epitaxially in the direction approximately perpendicular to the solidification interface where the temperature gradient is the highest. Therefore, the columnar dendrites are inclined to the laser scanning direction. As shown in Fig. 3b, under the low linear heat input, QCW LPD sample exhibits refined equiaxed dendritic microstructures. Through the comparison of Fig. 3a, b, it is obvious that the QCW laser mode can refine the grain size efficiently and eliminate the directivity of crystal orientation. This is because for the CW mode, the transient temperature of the molten pool is relative stable, while for the QCW mode, the transient temperature of molten pool exhibits periodical oscillations due to the periodical laser energy input. During the laser pulse intervals under the QCW laser mode, the temperature rapidly increases over the melting point of Inconel 718 when the laser is turned on, and then cools sharply below the melting

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Fig. 3 The microstructure in the laser-deposited Inconel 718 sample a 1, b 2, c 3 and d 4

point after the laser is turned off, resulting in a quenching effect on the molten pool and therefore increasing the nucleation rate and nuclei number. The periodical expansion and contraction of the molten pool not only change of the geometry of solidification interface and associated direction of the maximum thermal gradient but also affect the scour of liquid flow on the crystal tip, suppressing the continuous directional growth of the columnar dendrites. All these factors contribute to the formation of fine equiaxed dendrites and refinement of the solidification microstructure in the QCW sample. As shown in Fig. 3c, under the high linear heat input (120 J/mm), the crystallographic orientation of coarse columnar dendrites in the CW LPD sample tends to chaotic instead of inclination to the laser scanning direction. Comparing Fig. 3a, c,

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Fig. 4 Tensile strength of laser-deposited Inconel 718 at 600 °C

for the CW laser mode, it is evident that the increase of linear heat input not only weakens the crystallographic orientation uniformity of coarse columnar dendrites in the laser-deposited sample but also enlarges the grain size. As shown in Fig. 3d, under the high linear heat input, the microstructure morphology in the QCW LPD sample shows equiaxed dendritic microstructures. Comparing Fig. 3b, d, for the QCW laser mode, the results indicate that the increase of linear heat input changes the microstructure morphology little while enlarging grains size in the QCW LPD sample. This is because the high linear heat input increases the temperature of molten pool, which reduces the cooling rate and prolongs the growth time of crystal. At the same time, the high linear heat input weakens the quenching effect of the QCW laser mode, resulting in some columnar crystals form among the equiaxed grains. Figure 4 shows the tensile properties of laser-deposited Inconel 718 measured at 600 °C. Under the low linear heat input, the tensile strengths of QCW LPD and CW LPD are 1255 MPa and 1185 MPa, respectively. The tensile strength of QCW LPD sample is 6% larger than that of CW LPD sample. Under the high linear heat input, the tensile strengths of QCW LPD and CW LPD are 1169 MPa and 1066 MPa, respectively. The tensile strength of QCW LPD is 9.6% higher than that of CW LPD. The results suggest that low linear heat input advances tensile strength. Under the same condition, the QCW LPD can obtain higher tensile strength than CW LPD. According to the Hall–Petch equation, the smaller the grain size is, the better the mechanical properties are. As shown in Fig. 3, QCW laser mode and low linear heat input benefit the refinement of grain size, while CW laser mode and high heat input enlarge the grain size. Therefore, in this study, sample with the low linear heat input (60 J/mm) and QCW laser mode obtains the highest tensile strength 1255 MPa, sample with the high linear heat input (120 J/mm) and CW laser mode gets the lowest tensile strength 1066 MPa. Due to the finer grain size, the tensile strength at 600 °C of sample with low linear heat input (60 J/mm) and QCW laser mode is 17.7% higher than that with high linear heat input (120 J/mm) and CW laser mode. As shown in Fig. 5a, under the low linear heat input, the fracture surface of the QCW LPD sample is featured by deep equiaxed dimples. Therefore, the fracture mechanism is a transgranular ductile mode. As shown in Fig. 5b, the fracture surface of the CW LPD sample is featured by dendritic fracture patterns, indicating that the

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Fig. 5 Fracture surface of a QCW LPD and b CW LPD Inconel 718 samples with low linear heat input (60 J/mm)

fracture occurred preferentially along interdendritic regions. The deeper and larger dimples in the fractured QCW LPD sample (Fig. 5a) indicate a higher ductility of the QCW LPD sample than the CW sample (Fig. 5b). The fine and equiaxed dendritic microstructure contributes to the improved tensile property of the QCW LPD sample.

Conclusions In this study, the microtexture and tensile properties of CW LPD and QCW laser powder-deposited (LPD) Inconel 718 were studied under different linear heat inputs. The results showed that under low linear heat input, the microtexture morphologies with CW LPD and QCW LPD were long directional columnar grains and chaotic grains, respectively. The grain size of QCW LPD is finer than that of CW LPD. With the increase of linear heat input, the directivity of microtexture morphology with CW LPD becomes unobvious, while the microtexture morphology of QCW LPD changes little with enlarged grain size. Due to the finer grain size, the hightemperature tensile strength of sample with low linear heat input (60 J/mm) and QCW laser mode is 17.7% higher than that with high linear heat input (120 J/mm) and CW laser mode. Acknowledgements This work was supported by the Shenzhen Science and Technology Innovation Commission (Grant No. JCYJ20170817111811303) and Shenzhen Key Laboratory for Additive Manufacturing of High-Performance Materials (Grant No. ZDSYS201703031748354).

References 1. Qi H, Azer M, Singh P (2010) Adaptive toolpath deposition method for laser net shape manufacturing and repair of turbine compressor airfoils. Int J Adv Manuf Technol 48(1–4):121–131

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2. Mazumder J, Dutta D, Kikuchi N, Ghosh A (2000) Closed loop direct metal deposition: art to part. Opt Lasers Eng 34(4):397–414 3. Wu X, Mei J (2003) Near net shape manufacturing of components using direct laser fabrication technology. J Mater Process Technol 135(2–3):266–270 4. Sexton L, Lavin S, Byrne G, Kennedy A (2002) Laser cladding of aerospace materials. J Mater Process Technol 122(1):63–68 5. Lu Z, Li D, Tong Z, Lu Q, Traore M, Zhang A, Lu B (2011) Investigation into the direct laser forming process of steam turbine blade. Opt Lasers Eng 49(9):1101–1110 6. Lu Z, Zhang A, Tong Z, Yang X, Li D, Lu B (2011) Fabricating the steam turbine blade by direct laser forming. Mater Manuf Process 26(7):879–885 7. Gu D, Meiners W, Wissenbach K, Poprawe R (2012) Laser additive manufacturing of metallic components: materials, processes and mechanisms. Int Mater Rev 57(3):133–164 8. Mercer C, Soboyejo ABO, Soboyejo WO (1999) Micromechanisms of fatigue crack growth in a single crystal Inconel 718 nickel-based superalloy. Acta Mater 47(9):2727–2740 9. Zhao X, Chen J, Lin X, Huang W (2008) Study on microstructure and mechanical properties of laser rapid forming Inconel 718. Mater Sci Eng A 478(s 1–2):119–124 10. Trosch T, Strößner J, Völkl R, Glatzel U (2016) Microstructure and mechanical properties of selective laser melted Inconel 718 compared to forging and casting. Mater Lett 164:428–431 11. Engler O, Randle V (2009) Introduction to texture analysis: macrotexture, microtexture, and orientation mapping. CRC press 12. Blackwell P (2005) The mechanical and microstructural characteristics of laser-deposited IN718. J Mater Process Technol 170(1):240–246 13. Ocelík V, Furár I, De Hosson JTM (2010) Microstructure and properties of laser clad coatings studied by orientation imaging microscopy. Acta Mater 58(20):6763–6772 14. Moat R, Pinkerton A, Li L, Withers P, Preuss M (2011) Residual stresses in laser direct metal deposited Waspaloy. Mater Sci Eng A 528(6):2288–2298 15. Liu F, Lin X, Huang C, Song M, Yang G, Chen J, Huang W (2011) The effect of laser scanning path on microstructures and mechanical properties of laser solid formed nickel-base superalloy Inconel 718. J Alloy Compd 509(13):4505–4509 16. Lambarri J, Leunda J, García Navas V, Soriano C, Sanz C (2013) Microstructural and tensile characterization of Inconel 718 laser coatings for aeronautic components. Opt Lasers Eng 51(7):813–821

Part VIII

Additive Manufacturing: Materials Design and Alloy Development

Alloy Design for Biomedical Applications in Additive Manufacturing K.-P. Hoyer and M. Schaper

Abstract Since bioresorbable metal alloys like magnesium and iron are highly interesting in biomedical applications, significant efforts have been made to decrease the degradation rate of magnesium alloys, as well as to increase the degradation rate of iron-based alloys. Since silver is known to act as an effective cathode for iron, in addition to its antibacterial behaviour, it is gaining more interest as an alloying element for iron. Unfortunately, silver itself corrodes extremely slowly, and the effect of silver in the nanoscale on the human organism is, up to now, still controversial. Therefore, new silver alloys based on rare earth elements, as well as on typical elements for biomedical applications (Ca, Mg, Zn) with an adapted degradation profile (rate, scale of corrosion products), are a focal point of this work. Due to the immiscibility of iron and silver, it is not possible to cast iron–silver X-alloys, but it is feasible to manufacture these alloys using powder-based additive technologies. Keywords Iron-based alloys · Biomedical applications · Alloy design · Selective laser melting

Introduction Iron-based alloys are in focus and are still gaining increasing attention as bioresorbable materials for biomedical implants, such as stents or osteosynthesis plates [1–8]. Due to their slow degradation rate and high strength, iron-based alloys have to be modified to achieve adapted degradation rates and mechanical properties matching those of human bones. A promising approach is to alloy silver. Silver itself is biocompatible and antibacterial [9, 10] and should increase the degradation rate due to its positive electrochemical potential [11–13]. Huang et al. showed that silver and K.-P. Hoyer (B) · M. Schaper Chair of Materials Science, Paderborn University, Paderborn, Germany e-mail: [email protected] K.-P. Hoyer · M. Schaper Direct Manufacturing Research Center, Paderborn University, Paderborn, Germany © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_44

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gold are effective cathodes, which increase the corrosion rate of iron-based alloys [11]. However, in terms of biomedical applications, silver release and subsequent silver degradation have to be addressed as an enrichment of silver in a biological environment cannot be tolerated. Since the particles released during the degradation of the implant in the tissue are to be phagocytized, they must consist of a modified, biocompatible, non-corrosion-resistant noble metal alloy still to be developed. A promising candidate is calcium, as calcium is an endogenous element in the human body. Unfortunately, it is not possible to cast iron–silver alloys. Iron and silver are total immiscible in the solid as well as in the liquid state [14]. New manufacturing technologies, which allow the processing of conventionally immiscible material systems, must be considered as a manufacturing pathway. One such pathway is additive manufacturing. With additive manufacturing, it is possible to overcome these disadvantages and, in addition, to allow unique processing conditions, like unprecedented design freedom and individualized geometries [15–17]. In previous research, we have already proved the processability of an iron–manganese–silver alloy by applying selective laser melting [12, 13]. The mechanical properties, as well as the corrosion behaviour of the prepared alloys are promising. Therefore, the first attempt to modify the corrosion behaviour of the silver component by alloying calcium is addressed within this research.

Experimental Materials The materials investigated are two silver–calcium alloys with different calcium contents, as well as pure silver as a reference. Silver pallets (ESG Edelmetall-Service GmbH & Co. KG, Germany) and calcium pallets (99.5%, HMW Hauner GmbH & Co. KG) were used as base plate materials. All alloys were cast under argon atmosphere into small graphite crucibles and cooled down slowly. Afterward, small plates (1 × 15 × 10 mm3 ) were prepared by electrical discharging. Before starting the microstructural investigations and the analysis of the corrosion behaviour, all specimens were grinded, polished and cleaned with ethanol. In total, two different calcium contents have been investigated throughout this study: one alloy containing 6.0 wt% Ca and another system containing 10.0 wt% Ca. These calcium contents either lead to a eutectic microstructure or to defined intermetallic phases, as well as a lower melting point (Fig. 1). Table 1 summarizes the theoretical as well as the nominal chemical composition measured by means of energy dispersive spectroscopy (EDS), for all specimens. According to the analysis, the nominal composition matches the theoretical composition quite good.

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Fig. 1 Binary phase diagram of the AgCa system according to [14] Table 1 Theoretical and nominal composition of the addressed alloy systems; the nominal composition was measured by means of energy dispersive X-ray spectroscopy (EDS)

Alloy

Pure Ag

AgCa6

AgCa10

Theoretical composition

wt% Ag

100

94.00

90.00

wt% Ca

0

6.00

10.00

Nominal composition

wt% Ag

100

93.84

90.13

wt% Ca

0

6.16

9.87

Methods The immersion tests, as well as the open-circuit potential (OCP), were measured in Ringer’s lactate solution (131.00 mmol/l sodium, 5.40 mmol/l potassium, 1.80 mmol/l calcium, 112.00 mmol/l chloride, 28.00 mmol/l lactate) for 1200 h (immersion test) or, respectively, 600 s (OCP). OCP measurements were carried out using an MLab 100 potentiostat from Bank Elektronik with a scan rate of 1 scan/s for 600 s. For the immersion tests, a setup with beakers, in which the Ringer’s lactate solution was circulated with a small pump, was employed. The beakers were filled with 100 ml solution, and one specimen was suspended in each beaker. To suspend the specimens, small holes with a diameter of 0.7 mm were drilled into one corner of each of the specimens, and a nylon thread was pulled through this hole. The beakers were then

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Table 2 Composition of the etching solutions and etching parameters Chemicals

Ratio

Etching parameters

Etching solution 1

Aqueous ammonia solution (25%)

Aqueous hydrogen peroxide solution (35%)

5:1

3 × 3 s for Ag, AgCa6; 3 × 2 s for AgCa10

Etching solution 2

Deionised water

Aqueous hydrochloric acid (37%)

10:1

3×5s

covered to avoid any contamination of the test solution. Small, flexible tubes helped to flush the test solution with air to avoid a depletion of oxygen and to realize a smooth moving of the solution. Possible corrosion products on the specimens were removed carefully, and the test conditions were kept constant throughout the duration of the immersion tests. Every week, the test solution was removed and replaced by a freshly mixed solution. Meanwhile, every specimen was removed for a short period, dried and weighted (XP205 DeltaRange, Mettler Toledo) to get immediate information about the degradation rate. After 1200 h, all specimens were removed from the test solution, the corrosion products were carefully removed, and the specimens were cleaned with ethanol, followed by a subsequent ultrasonic cleaning in an ethanol bath for 5 min. Next, all specimens were rinsed with ethanol and stored in a desiccator until they were weighted and prepared for the microstructural investigations. To investigate the microstructure, optical light microscopy (OM/LM, Keyence VHX 5000, Olympus SZX or Zeiss Axiophoto) and scanning electron microscopy (SEM, Zeiss Ultra Plus) were employed before and after the corrosion testing, respectively. For this, all specimens were etched alternating in two etching solutions according to Table 2. Further analysis regarding the chemical composition was performed with energy dispersive X-ray spectroscopy (EDS).

Results and Discussion Corrosion Analysis Alloying silver with calcium should lead to a faster degradation of the binary alloy compared to pure silver. Since the open-circuit potential is an indicator for the susceptibility to corrode, the open-circuit potential of the two AgCa alloys was measured in Ringer’s lactate solution for 600 s. Figure 2 depicts the results of these measurements. It can be seen that the OCP of both alloys is more negative as compared to pure silver, which is equivalent to a higher susceptibility to corrode. A higher amount of calcium shifts the OCP to a more negative value as calcium is a less noble material

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Fig. 2 Open-circuit potential of the two AgCa alloys compared to the reference (pure silver) in Ringer’s lactate solution

than silver. The average OCPs are as follows: 255 mV (Ag), 129 mV (AgCa6), and 92 mV (AgCa10), with pure silver showing a less noble behavior than in standard conditions (e.g., RT, 1025 mbar, standard test solution). These results are supported by the results from the immersion tests (Fig. 3). Pure silver hardly corrodes in the test solution, whereas both alloy systems show a significant increase in their mass losses throughout the exposure time of 1200 h. In accordance with the OCP, a higher amount of calcium leads to an increased degradation rate. The average mass losses for the two AgCa alloys after 1200 h are as follows: 17.31 mg/cm2 (AgCa6) and 28.54 mg/cm2 (AgCa10).

Microstructural Analysis The etched specimens were grinded and polished for the microstructural investigations as mentioned above. To obtain the microstructure of the different alloys, light microscopy was employed (Figs. 4 and 5). In comparison with the AgCa10 alloy, the AgCa6 alloy (Fig. 4) shows a completely different microstructure. In accordance with the binary phase diagram (Fig. 1), AgCa6 reveals a eutectic microstructure. Therefore, the degradation behavior of AgCa6 should be more homogenous than the degradation of the AgCa10 alloy. In contrast, the microstructure of the AgCa10 alloy (Fig. 5) shows intermetallic phases, which are embedded in the eutectic matrix of the alloy. According to the above shown binary phase diagram, these phases should be either Ag7 Ca2 or Ag9 Ca2 (Fig. 1). Due

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Fig. 3 Average mass losses of the Ag, AgCa6 and AgCa10 alloys in Ringer’s lactate solution during the immersion tests

Fig. 4 Cross sections of the AgCa6 alloy

Fig. 5 Cross sections of the AgCa10 alloy

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Fig. 6 EDS mappings of the AgCa alloys showing the homogenous distribution of calcium (dark spots) in the silver matrix (red), a AgCa6, b AgCa10

to the occurrence of intermetallic phases, the degradation rate of this alloy should be faster than that of AgCa6. In general, intermetallic phases are known to have a more positive OCP than the matrix material (silver) based on their metallic–ceramic behavior. Due to this, these resulting small galvanic cells significantly increase the degradation rate of the silver matrix and, therefore, of the alloy. Consequently, the OCP of this alloy was shifted to a more negative value (Fig. 2). EDS mappings of the casted AgCa alloys verified the homogenous distribution of calcium in the silver matrix (Fig. 6). Further surface investigations with light microscopy after 600 h and 1200 h showed the influence of the electrolyte on the alloys. For the AgCa6 alloy, the surface showed no preferred areas with corrosion-induced mass loss (Fig. 7). The entire surface corrodes more or less evenly without any formation or breakdown of a protective film or the occurrence of selective corrosion or contact corrosion. In contrast, the AgCa10 alloy shows a local corrosion attack between the eutectic microstructure and the intermetallic phases (Fig. 8). Contrary to expectation, this behavior is more pronounced after 600 h immersion time than after 1200 h. Typically, the corrosion attack would be more intense with increasing exposure time in corrosive media, if the specimen does not form any protective layer or if corrosion products cover the entire specimen surface. Even if a layer on the surface was formed, localized corrosion would occur more often with increasing exposure time. Therefore, it is quite untypical that localized corrosion is more pronounced after a shorter exposure time.

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Fig. 7 Cross sections of AgCa6 after 600 h (a, c) and after 1200 h (b, d) immersion in Ringer’s lactate solution

Conclusion and Outlook The microstructural, as well as electrochemical and immersion, test results prove that alloying silver with calcium leads to a significant increase in the degradation rate of the binary alloy system. Therefore, the addition of calcium to silver is a promising approach to adapt the degradation rate of iron-based silver used as material for biomedical implants. In further investigations, the size of the particle release is examined to verify, if these can be phagocytized. If this can be proven, the FeMnAgCa alloy would be casted in a bigger ingot and gas-atomized for applications in the area of selective laser melting.

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Fig. 8 Cross sections of AgCa10 after 600 h (a, c) and after 1200 h (b, d) immersion in Ringer’s lactate solution

Acknowledgements The authors thank Mathias von Spalden, Maik Robert Peters, and Thomas Janzen for their support during the investigations.

References 1. Peuster M, Hesse C, Schloo T, Fink C, Beerbaum P, von Schnakenburg C (2006) Long-term biocompatibility of a corrodible peripheral iron stent in the porcine descending aorta. Biomaterials 27:4955–4962 2. Hermawan H, Moravej M, Dubé D, Fiset M, Mantovani D (2007) Degradation behaviour of metallic biomaterials for degradable stents. Adv Mater Res 15–17:113–118 3. Hermawan H, Dubé D, Mantovani D (2007) Development of degradable Fe-35Mn alloy for biomedical application. Adv Mater Res 15–17:107–112 4. Hermawan H, Purnama A, Dube D, Couet J, Mantovani D (2010) Fe-Mn alloys for metallic biodegradable stents: degradation and cell viability studies. Acta Biomater 6:1852–1860 5. Hermawan H, Dube D, Mantovani D (2010) Developments in metallic biodegradable stents. Acta Biomater 6:1693–1697

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6. Heiden M, Walker E, Nauman E, Stanciu L (2015) Evolution of novel bioresorbable iron— manganese implant surfaces and their degradation behaviors in vitro. J Biomed Mater Res Part A 103A:185–193 7. Drynda A, Hassel T, Fr-W Bach, Peuster M (2015) In vitro and in vivo corrosion properties of new iron–manganese alloys designed for cardiovascular applications. J Biomed Mater Res B 03:649–660 8. Francis A, Yang Y, Virtanen S, Boccaccini AR (2015) Iron and iron-based alloys for temporary cardiovascular applications. J Mater Sci Mater Med 26:138 9. Jansen B, Rinck M, Wolbring P, Strohmeier A, Jahns T (1994) In vitro evaluation of the antimicrobial efficacy and biocompatibility of a silver-coated central venous catheter. J Biomater Appl 9:55–70 10. Karchmer TB, Giannetta ET, Muto CA, Strain BA, Farr BM (2000) Randomized crossover study of silver-coated urinary catheters in hospitalized patients. Arch Intern Med 21:3294–3298 11. Huang T, Cheng J, Bian D, Zheng Y (2016) Fe-Au and Fe-Ag composites as candidates for biodegradable stent materials. J Biomed Mater Res B 104(2):225 12. Niendorf T, Brenne F, Hoyer P, Schwarze D, Schaper M, Grothe R, Wiesener M, Grundmeier G, Maier HJ (2015) Processing of new materials by additive manufacturing: iron-based alloys containing silver for biomedical applications. Metall Mater Trans A 46:2829–2933 13. Wiesener M, Peters K, Taube A, Keller A, Hoyer K-P, Niendorf T, Grundmeier G (2017) Corrosion properties of bioresorbable FeMn-Ag alloys prepared by selective laser melting. Mater Corros 68:1028–1036 14. ASM Alloy P D D ASM alloy phase diagram database 15. Murr LE, Gaytan SM, Ramirez DA, Martinez E, Hernandez J, Amato KN, Shindo PW, Medina FR, Wicker RB (2012) Metal fabrication by additive manufacturing using laser and electron beam melting technologies. J Mater Sci Technol 28:1–14 16. Yan C, Hao L, Hussein A, Raymont D (2012) Evaluations of cellular lattice structures manufactured using selective laser melting. Int J Mach Tools Manuf 62:32–38 17. Nune KC, Devesh R, Misra K, Gaytan SM, Murr LE (2014) Interplay between cellular activity and three-dimensional scaffold-cell constructs with different foam structure processed by electron beam melting. J Biomed Mater Res 103A(5):1677–1692

Surface Inoculation of Aluminium Powders for Additive Manufacturing Guided by Differential Fast Scanning Calorimetry Lennart Tasche, Kay-Peter Hoyer, Evgeny Zhuravlev, Guido Grundmeier, Mirko Schaper and Olaf Keßler Abstract To improve the solidification process of laser beam melting (LBM), and thus enable the printing of hard-to-weld high-strength aluminum alloys, the project is designed to modify the powder surface by adding nanoparticles for a guided nucleation. In situ testing of the fast melting and solidification process of single particles will be performed using differential fast scanning calorimetry (DFSC). DFSC results will be transferred to LBM of aluminum alloys and used to adjust parameter settings. Specimens will be printed, to check crack formation and physical as well as mechanical material properties. The aluminum alloy EN AW-7075 (EN AW-AlZnMgCu1.5, Al7075) is the most common high-strength aluminum alloy used in mechanical engineering, aircraft industry and, due to new developments like press quenching, also in the automobile industry. Due to the importance of Al7075, this research focusses on inoculating its powder particles and thus, enables crack and porosity free 3D printing. Keywords High-strength alloys · Laser beam melting · Surface inoculation · Grain refinement · DSC

L. Tasche (B) · K.-P. Hoyer · M. Schaper Chair of Materials Science, Paderborn University, Warburgerstrasse 100, 33098 Paderborn, Germany e-mail: [email protected] E. Zhuravlev Institute of Physics and °CALOR, University of Rostock, Albert-Einstein-Str. 25, 18051 Rostock, Germany G. Grundmeier Technical and Macromolecular Chemistry, Paderborn University, Warburger Str. 100, 33098 Paderborn, Germany O. Keßler Chair of Materials Science, University of Rostock, Albert-Einstein-Str. 2, 18059 Rostock, Germany L. Tasche · K.-P. Hoyer · G. Grundmeier · M. Schaper DMRC, Paderborn University, Mersingweg 3, 33100 Paderborn, Germany © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_45

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Introduction Additive manufacturing of high-strength aluminum and steel alloys is still restricted. Crack formations and porosity occurring during the solidification process limit the available number of printable alloys. To increase the number and enable the printing of hard-to-weld alloys, the improvements in aluminum welding shall be introduced into additive manufacturing [1, 2]. With the high thermal diffusivity of aluminum, up to 6 times greater than for steel, a large area of the workpiece is heat affected. In additive manufacturing, the heat-affected area increases, due to the repeated heating of the layers above and the overlapping melt pool. The continuous reheating leads to a release of dislocations and grain enlargement, which has a negative impact on the mechanical properties [1, 3]. Furthermore, a wide solidification temperature range, as in high-strength aluminum alloys, leads to dendritic growth of the grains and thus, to hot tear defects. Due to the dendritic growth, hollow areas are created during the solidification shrinkage [4]. Therefore, grain refinement techniques are used, to control the microstructure of the heat-treated area, and thus, in additive manufacturing, of the whole workpiece. Figure 1 shows scanning electron microscope (SEM) images of two parameter studies, carried out to deepen the knowledge regarding the laser beam melting (LBM) process of Al7075. The adjusted parameters have been the energy density (Fig. 1a–c) and the preheating of the building platform (Fig. 1d–f). With adjusted parameters, the build quality can be increased but, due to the solidification shrinkage, porosity and especially hot cracking are not fully preventable. Methods and mechanisms of grain refinement have been extensively studied, e.g. by [5–9]. The inoculation of the alloy by adding heterogeneous nuclei is a common method of grain refinement, which was implemented in the additive manufacturing process of Al7075 during this project. This allows to reduce the interface energy and allows aluminum grains to grow on the nuclei surface. Typical grain refiners for Al7xxx are TiB2 and TiC, which will be used in this project [7].

Objectives To extend the areas of application of 3D-printed parts, it is necessary to investigate the accurate usage of grain refiners in the LBM process. The objective is to achieve crackfree printing of 7xxx aluminum using the accurate type and amount of nanoparticles. Furthermore, parameter studies need to be carried out for the particle surface addition, as well as the LBM of the modified powder. Four research hypotheses have been compiled in the project: (i) Grain refinement leads to an improvement of the solidification conditions in the molten pool and thus enables crack-free printing of aluminum alloys. (ii) Grain refinement is achievable by adding nanoparticles to the metal powder surface (this modification can be performed on large powder volumes). (iii) The efficiency of the grain refinement on single

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Fig. 1 Al7075 LBM parameter studies on cracks and porosity. a–c Variation in energy density, d–f variation in preheating of building platform

particles can be in situ determined using the differential fast scanning calorimetry method. (iv) The mechanical properties as well as the corrosion resistance of the produced alloy can be correlated to the grain refinement.

Work Program The work program is divided into steps that are carried out by the four partners of the project group. After a quick introduction to the steps, selected steps and methods will be discussed in detail. After the selection of the nuclei, the surface inoculation of the Al 7xxx powder particles is performed by the group of Technical and Macromolecular Chemistry of Paderborn University. Afterwards, using an advanced differential scanning calorimetry, the differential fast scanning calorimetry (DFSC), the competence center °CALOR of the University of Rostock investigates single modified powder particles, by subjecting them to multiple melting and solidification cycles for in situ testing. Using single modified powder particles, an immediate feedback to the modification can be provided. After successful testing, bigger amounts of the powder are produced, to allow LBM testing of defined test specimens and thus enable material characterization by the group of Materials Science of the Paderborn University.

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Fig. 2 Schematic drawing of the surface inoculation process

Powder Surface Inoculation Different methods have been tried for the powder surface inoculation. The method implemented in this project is a dispersion/drying-on process. This allows a successful adhesion as well as an even distribution of large amounts of nuclei and is furthermore scalable to large amounts of powder. The nanoparticles are suspended in an aqueous electrolyte dispersion containing polyethylenimine (PEI). To ensure proper resolving, ultrasound is applied. Remaining agglomerates sediment and the dispersion can be applied to the aluminum powder. The method leads to a fast and homogenous wetting of the origin powder. The drying process is supported by heating and pumping. The chemical equation is shown in Eq. 1, and the chemical procedure is shown in Fig. 2. Al2 O3 + 2O H − + 3H2 O → 2 AL(O H )− 4

(1)

The initial state in Fig. 2 is an Al alloy particle, covered in a passive film and wetted. The nuclei nanoparticles are covered in a thin electrolyte, containing PEI to adjust the surface charge. During the water evaporation, the pH increases. After the evaporation, the nanoparticles are strongly absorbed on the Al particles surface. The process leads to a modified surface oxide film. The etching of the native Al oxide film due to the intermediate alkaline pH is indicated in Fig. 3d. First results of inoculated Al 7075 are shown in Fig. 3. Figure 3a, b shows SEM images of the unmodified powder. Figure 3c, d shows the same powder, modified with 50 nm titanium carbide (TiC) nanoparticles. Figure 3f, g shows EDX mappings of the inoculated powder. Figure 3e shows the SEM image as a reference area. In Fig. 3f, the aluminum content is measured, and the titanium content is measured in Fig. 3g. An even distribution of nanoparticles is visible. Furthermore, changes in the particles surface are apparent. Besides the previously discussed change in the particles oxygen layer, colorimeter tests show a 30% increase

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Fig. 3 SEM images of TiC inoculated Al7075 powder: a, b unmodified, c, d modified, e–g EDX mapping

in the black level of the particle, which indicates a higher absorption rate and may result in a decrease in the required laser beam energy.

DFSC Differential fast scanning calorimetry has been developed by the competence center °CALOR as an advanced version of differential scanning calorimetry. Ten micronsized samples enable high heating and cooling rates in an accordingly small measuring cell. Figure 4 shows the composition of the calorimeter, with an overview of the measuring cell in Fig. 4a and a cross section in Fig. 4b. Heaters and thermopiles with a smaller heat capacity, made out of doped polysilicon, are deposited on the 500-nm-thin silicon nitride membrane [10–12].

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Fig. 4 a DFSC measuring cell overview, b schematic cross section

DFSC has further been extended by the differential reheating method (DRM). Thus, exothermal precipitation reactions during cooling can be detected as an equivalent endothermal dissolution reaction during subsequent reheating. This method can be applied at rates of several 103 K/s, allowing measurements in the high-sensitivity range of DFSC. Among others, this development has been used to detect precipitation reactions during rapid quenching of aluminum alloy 7150. A further development, extreme quenching with rates of 105 K/s, allows solidification studies under LBM conditions [10]. Thereby, this method enables in situ analyses of single inoculated powder particles and a dynamic and fast feedback about the powder characteristics, compared to the inoculation of big amounts of powder for test prints. Figure 5a shows the melting and solidification traces of AlSi10 Mg for different heating and cooling rates. Figure 5b shows transmission electron microscopy (TEM) images of the aluminum particle after DFSC measurements at 5000 K/s. Due to the high cooling and heating rates, the DFSC method is perfectly suited for the simulation of the LBM process. In addition, the melting and solidification procedure can be iterated to describe the actual process during LBM, especially of hard-to-weld aluminum alloys, with the previously discussed overlapping melt pools and large heat-affected areas. Current investigations assume an effect on the nuclei and the grain refinement, dependent on the prior heat treatment. The influence of the grain refinement on the undercooling during crystallization is shown in Fig. 6a, b. The effect of smaller grains leading to a smaller temperature difference θ and a shorter undercooling period enables a first assumption of the achieved grain refinement from the DFSC results. In addition, the tested single powder particles are prepared, using a focused ion beam (FIB) system, and then analyzed, using SEM and TEM, to determine the microstructure after solidification. The small amount of particles needed allows a large number of tested nuclei and inoculation conditions. The results are immediately reported to the previous work packages and

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Fig. 5 a DFSC traces of a single AlSi10 Mg particle during heating and cooling. b TEM image of the particle after DFSC measurements at 5000 K/s

Fig. 6 a Thermal analysis during nucleation (schematically), b grain size depending on undercooling and period of undercooling (schematically) [7]

allow a targeted modification of the inoculation process. Promising parameters are then reproduced in a larger scale for LBM testing.

LBM Specimens for microstructural analysis and mechanical testing are manufactured on an SLM® 280HL with a 280 × 280 × 350 mm3 build chamber and a 400 W Yttrium laser. Proceeding from standard parameters for Al7075, in excerpts shown in Fig. 1, parameter studies are carried out on the manipulated powder to reach the full potential of the alloy. Printed specimens are 10 × 10 × 10 mm3 cube specimen for porosity and crack analysis. Grain sizes are measured using SEM and energy

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dispersive X-ray spectroscopy (EDX), with a successful nucleation implying that nanoparticles operate epitaxial as nucleation agents. Further, specimens for tensile and fatigue testing are produced. Processed material is used for corrosion testing by the Department of Technical and Molecular Chemistry. Passivation and local corrosion attacks are expected due to the powder modification. Potentiodynamic polarization experiments are carried out, using aerated 0.1 NaCl with a pH from 4–5 and 10–11, to analyze the corrosion properties. The results of the microstructural analysis and the mechanical and corrosion testing are used to adjust the parameter settings of the LBM process as well as for a feedback for the previous steps such as the inoculation process.

Conclusion During this project, a powder surface inoculation method shall be developed, which allows quick and successful inoculation of additive manufacturing Al7075 powder with nanoparticles functioning as nuclei in the melt. This method allows to induce grain refinement and thus enables the laser beam melting of hard-to-weld, highstrength aluminum alloys. In combination with the DFSC, small amounts of inoculated powder can be used to determine the effect of the chosen nanoparticle and the process as well as its parameters during the inoculation. The gained information can be fed to the previous work packages instantly and enable far-reaching study on the nanoparticles and the process.

References 1. Schulze G (2010) Die Metallurgie des Schweißens. Springer, pp 531–532. ISBN 978-3-64203182-3 2. Hu B, Richardson IM (2006) Mechanism and possible solution for transverse solidification cracking in laser welding of high strength aluminium alloys. Mater Sci Eng A 429:287–294 3. Dilthey U (2005) Schweißtechnische Fertigungsverfahren 2. Springer, pp 217–237. ISBN 9783-540-21674-2 4. Pabel T, Bozorgi S, Kneissl C, Faerber K, Schumacher P (2012) Einfluss der Legierungselemente auf die Heißrissneigung bei AlSi7MgCu-Gusslegierungen. Giesserei 99:30–37 5. Greer AL, Cooper PS, Meredith MW, Schneider W, Schumacher P, Spittle JA, Tronche A (2003) Grain refinement of aluminium alloys by inoculation. Adv Eng Mater 5(1-2):81–91 6. Quested TE (2004) Understanding mechanisms of grain refinement of aluminium alloys by inoculation. Mater Sci Technol 20(11):1357–1369; Surface inoculation of aluminium powders for additive manufacturing guided by differential fast scanning calorimetry P 22 of 31 7. Spittle JA (2006) Grain refinement in shape casting of aluminium alloys. Int J Cast Met Res 19(4):210–222 8. Greer AL (2016) Overview: application of heterogeneous nucleation in grain-refining of metals. J Chem Phys 145(21), art. no. 211704

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9. Guan RG, Tie D (2017) A review on grain refinement of aluminum alloys: progresses, challenges and prospects. Acta Metall Sin (Engl Lett) 30(5):409–432 10. Minakov AA, Schick C (2007) Ultrafast thermal processing and nanocalorimetry at heating and cooling rates up to 1 MK/s. Rev Sci Instrum 78(7), art. no. 073902 11. Minakov AA, Schick C (2016) Heat conduction in ultrafast thin-film nanocalorimetry. Thermochim Acta 640:42–51 12. Minakov AA, Schick C (2015) Dynamics of the temperature distribution in ultra-fast thin-film calorimeter sensors. Thermochim Acta 605:205–217

Mechanical Behavior and Microstructure of Porous Ti Using TiC as Reinforcement Shiyuan Liu, Jian Wang, Tengfei Lu, Guibao Qiu and Hao Cui

Abstract The porous Ti/TiC composite, where TiC behaviors as alloy reinforcement, with 50% porosity, was successfully prepared by powder metallurgy technology using acicular urea as space holder. The pore size in Ti foams is mainly in the range of 500–800 μm with average size of 600 μm, which is a little smaller than that of urea particles due to the pore-shrinking effect during the sintering process. On the other hand, the porosity and compressive strength of the prepared Ti products in this work increase as the adding content of TiC reinforcement, from 50.61% and 68.93 MPa for 0% TiC, to 53.08% and 106.32 MPa for 8% TiC, respectively. Results indicated that through adding TiC reinforcement into porous Ti, the pore structure, namely, the microstructure, in line with the mechanical behaviors of Ti foams could be dramatically affected, with finer pore structure and enhanced mechanical property obtained in the final prepared Ti foams. Keywords Porous Ti · TiC reinforcement · Mechanical property · Porosity

Introduction Ti metal has been widely applied in the fields of aerospace, automotive and biomedical industries due to its unique features of low density, high specific strength, excellent chemical resistance and good biocompatibility [1, 2]. At the same time, porous materials have aroused extensive focus of attention in recent years based on their specific applications such as bone transplantation, sound adsorption and energy consumption [3, 4]. Given these, porous Ti being provided with combined advantages of Ti alloys S. Liu · J. Wang · T. Lu · G. Qiu (B) College of Materials Science and Engineering, Chongqing University, Chongqing 400040, China e-mail: [email protected] H. Cui (B) State Key Laboratory of Power Transmission Equipment & System Security and New Technology, Chongqing University, Chongqing 400044, China e-mail: [email protected] © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_46

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and foams is emerging as a novel kind of functional material in the modern material engineering, especially winning significant influence and attraction in the biomedical field. However, this type of material inevitably suffers from the relatively lower strength weakened by the pores inner prepared specimens [5]. Apart from that, the strength and stiffness of the materials would drop accordingly as their increasing porosity, and the plateau region of stress–strain curve during their compressive tests might disappear owing to the increased ratio of the closed cellular pores [1, 6]. Recently, the fabrication of Ti matrix composites (TMC) is proposed as a promising manner to enhance the mechanical properties of Ti-based materials through the addition of certain metals or nonmetal components. Specifically, the compressive or tensile strengths of prepared TMC would be dramatically improved compared with those properties of pure Ti metal, especially when they are used in combination with the advantages of powder metallurgy techniques. This kind of foam, integrated advantages of both Ti and porous structure, has recently become one of the most attractive and influent materials in modern technology. The possibility of reinforcement in porous metal matrix composites can modify properties such as mechanical behavior of material so that their applications could be broadened; especially, they are required to be applied in some complex environment. Among reinforcements, titanium carbide (TiC) has the advantages of high stiffness and hardness, superior heat conduction performance, and above all excellent compatibility with Ti [7–10]. The use of TiC in Ti alloys for reinforcement has attracted much attention in recent years with the aim of improving their mechanical performance due to their density, elastic modulus, chemical and thermal stability being similar to those of Ti metal [11, 12]. In this study, a porous Ti/TiC composite was prepared using Ti as basis and TiC as reinforcement by powder metallurgy technology offering advantages such as low cost, near-net-shape fabrication, increase in material yield, and variation of composition. The high energy ball-milling technology was applied to preprocess Ti powder for improving its property. The mechanical property and microstructure of porous Ti matrix composites were studied.

Experimental Method The raw materials utilized for the preparation of porous Ti/TiC composites were Ti powder, TiC powder and acicular urea particles, where the TiC behaves as the composite reinforcement while the urea as the spacer holder. Both Ti and TiC powders have the purity larger than 99.5% with powder size of 300 meshes. The chemical composition by mass percentage of Ti powder could be seen as follows: Si ≤ 0.03%, Fe ≤ 0.05%, Cl ≤ 0.03%, C ≤ 0.02%, O ≤ 0.32%, N ≤ 0.01%, H ≤ 0.04%. At the same time, the acicular urea particle has the average size ranging from 0.8 to 1.1 mm with porosity larger than 99%. In this study, the content of urea adding into the green sample of porous Ti was designed as 50% aiming at obtaining corresponding porosity for the specimens. Adding amounts of TiC in green compacts are defined

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to be 0wt%, 2wt%, 4wt%, 6wt%, and 8wt%, respectively, indicating corresponding Ti amounts as 100wt%, 98wt%, 96wt% 94wt% and 92wt%, accordingly, which totally occupies 50% volume amount of the final products. The mixed Ti and TiC powders were blended in a planetary ball mill at a speed of 300 r/min for 30 min in order to improve surface energy and increase the homogeneity of mixture. It should be mentioned that the mechanical properties of final Ti-based materials would be pronouncedly enhanced if the raw powders were previously ball-milled according to authors’ previous experiment. Then, blended power and urea particle were mixed together for 10 min to ensure the homogenous distribution of blended powers on the urea particles. Moreover, a little (about 0.5 ml) alcohol and zinc stearate (about 0.5 g) were applied to ensure the fluency and lubrication between mixed raw materials and cylindrical die during the green sample compressing process. The mixed particle was thereafter uniaxially compressed into a green compact (d  15 mm, h  10 mm) under the pressure of 200 MPa, and 1 min lasting time at this pressure was required to enable the effective and uniform pressing for the green sample. Subsequently, the green samples were put into the sintering furnace for implementing the heating process. The samples were first heated up to 400 °C at a constant heating rate of 10 °C/min and held at that point for 40 min aiming at completely decomposing the urea; and second sintered at 1250 °C for 2 h (the sintering temperature, duration time and pressure mentioned above were demonstrated sufficiently in our previous studies). Additionally, given that Ti would be pretty active in the surroundings at high temperature, the high purity argon (≥99.99%) was poured into and pumped out of the sintering chamber during the whole heating process to achieve the atmosphere requirement and remove the potential trace oxygen. Finally, the samples were taken out from the furnace after they have been cooled to ambient temperature. The pore structure of Ti samples was determined by scanning electron microscope (SEM), while the porosity was calculated according to the following equation [13, 14]: P 1−

ρ∗ ρs

(1)

where ρ ∗ represents the density of prepared Ti foams, which could be measured by dividing the whole weight into the whole  volume of the sample; ρs denotes the  density of the solid Ti ρs  4.503 g/cm3 . Meanwhile, the compressive strength was measured by material testing machine.

Results and Discussion The porous Ti/TiC matrix composites with porosity of 50 ± 3.5% and average pore size of 600 μm were successfully prepared by powder metallurgy technology. At the same time, the matrix ratio, porosity, and compressive strength of relevant samples are described in Table 1.

498 Table 1 Matrix ratio, porosity, and compressive strength of porous Ti

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Sample

TiC content (%)

Porosity (%)

Compressive strength (MPa)

Ti/TiC

0

50.61 ±0.2

68.93 ±1.8

2

51.59 ±0.15

86.73 ±2.1

4

52.05 ±0.17

91.72 ±1.6

6

52.84 ±0.15 103.39 ±1.9

8

53.08 ±0.16 106.32 ±2.2

Pore Structure and Porosity The pore structure of porous Ti is shown in Fig. 1. It has been reported that pores in porous metals could be divided into two types [15]: the macropores obtained by the decomposition of urea particles and the micropores gained by partially sintering of Ti powders on the pore walls that frequently have the size about several micrometers. Macropores in Ti could also be divided into two categories: interconnected pores usually formed in high porosity Ti and unconnected ones usually existed in low porosity samples. The pore size in the porous Ti as the figures show spans in the range of 500–800 μm with the average size of 600 μm. The pore size in the samples is apparently smaller than the size of spacer added into the samples before; the account for this phenomenon could be attributed to the shrink of pores during the sintering process at high temperature. Therefore, the porosity of the final product is all larger than designed in the study, which mainly because the amount of micropores mentioned above is larger than shrink of macropores that both occurred during the sintering process. Figure 2 describes the porosity of samples as the variation in mass friction of TiC. It is obviously found out that the porosity increased as the content of TiC increased. The relationship between porosity (P) and TiC content (T ) in this study is, to put it precisely, found to obey an exponential relation as shown by the fitting line and Eq. (2). P  e3.92+0.01T −6.61T (R  0.99993) 2

Fig. 1 Pore structure for TiC/Ti porous composite

(2)

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Fig. 2 Relationship between porosity and mass fraction of TiC

Mechanical Behavior Analysis The compressive strain–stress curve of composites with different TiC mass fractions is shown in Fig. 3. It could be seen from the strain–stress curves that the compressive stress of porous Ti/TiC composites increased as the mass fraction of TiC increased, with the underline assumption that the mechanical property could be significantly improved using TiC as reinforcement. The reason for this result mainly attributes to the superior hardness of TiC and its excellent compatibility with Ti basis. It has been reported [16, 17] that the stress–strain curves exhibit three distinct regions: (i) linear elastic region, (ii) plateau region, and (iii) densified region. In plateau region (i.e. in region ii), stress oscillates around an average stress value and remains relatively flat even with increase in strain value (the average stress in this region is termed as plateau stress). In densified region (i.e. in region iii), stress increased drastically with slight increase in strain. The strain–stress curves in the present work could be understood in the similar way although it shows an ill-defined densified region owing to relatively large pore size. The strain–stress curve in Fig. 2 presents a relatively obvious linear elastic region and plateau region that coincided well to that of porous material. However, the average spacer size used in this paper was 0.8–1.1 mm resulting in main pore size of 600–800 μm in samples that are easily collapsed during the compression. The larger the pore size is, the easier the collapse tends to occur in porous Ti. Thus, the stress finally constantly decreased with the increasing strain. This conclusion is consistent with the previous study of Byounggab Lee.

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Fig. 3 Compressive curve of porous Ti with variation of TiC

Conclusions In this paper, the porous Ti/TiC composite, where TiC behaviors as alloy reinforcement, with 50% porosity, was successfully prepared by powder metallurgy technology using acicular urea as space holder. The pore size in Ti foams is mainly in the range of 500–800 μm with average size of 600 μm, which is a little smaller than that of urea particles due to the pore-shrinking effect during the sintering process. The porosity and compressive strength of the prepared Ti products in this work increase as the adding content of TiC reinforcement, from 50.61% and 68.93 MPa for 0% TiC to 53.08% and 106.32 MPa for 8% TiC, respectively. Our experimental would be meaningful to give a clear insight into the enhancement effect of porous Ti through adding TiC reinforcement, where finer pore structure and enhanced mechanical property could be obtained in the finally prepared Ti foams.

References 1. Li BQ, Wang CY, Lu X (2013) Effect of pore structure on the compressive property of porous Ti produced by powder metallurgy technique. Mater Des 50:613–619 2. Jiang G, Li Q, Wang C, Dong J, He G (2015) Fabrication of graded porous titanium–magnesium composite for load-bearing biomedical applications. Mater Des 67:354–359 3. Liu PS, Qing HB, Hou HL (2015) Primary investigation on sound absorption performance of highly porous titanium foams. Mater Des 85:275–281 4. Liu PS, Qing HB, Hou HL, Wang YQ, Zhang YL (2016) EMI shielding and thermal conductivity of a high porosity reticular titanium foam. Mater Des 92:823–828

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5. Jian X, Hao C, Guibao Q, Yang Y, Xuewei L (2015) Investigation on relationship between porosity and spacer content of titanium foams. Mater Des 88:132–137 6. Esen Z, Bor S¸ (2007) Processing of titanium foams using magnesium spacer particles. Scr Mater (5):341–344 7. Jiang J-Q, Lim T-S, Kim Y-J, Kim B-K, Chung H-S (1996) In situ formation of TiC-(Ti-6Al-4V) composites. Mater Sci Technol (4):362–365 8. Tong XC, Fang HS (1998) Al-TiC composites in situ-processed by ingot metallurgy and rapid solidification technology: part I. Microstructural evolution. Metall Mater Trans (3):875–891 9. Pingwei LW, Xiaonong Z, Di Z, Renjie W, Fang BJ (1999) Growth mechanism of in situ synthesis of TiC/Ti composite reinforcing. Acta Metall Sin (5):536–541 10. Radhakrishna Bhat BV, Subramanyam J, Bhanu Prasad VV (2002) Preparation of Ti-TiB-TiC & Ti-TiB composites by in-situ reaction hot pressing. Mater Sci Eng (1):126–130 11. Choy MT, Tang CY, Chen L, Wong CT, Tsui CP (2014) In vitro and in vivo performance of bioactive Ti6Al4V/TiC/HA implants fabricated by a rapid microwave sintering technique. Mater Sci Eng 42:746–756 12. Tang CY, Wong CT, Zhang LN, Choy MT, Chow TW, Chan KC, Yue TM, Chen Q (2013) In situ formation of Ti alloy/TiC porous composites by rapid microwave sintering of Ti6Al4V/MWCNTs powder. J Alloy Compd 557:67–72 13. Hu Y, Grainger DW, Winn SR, Hollinger JO (2002) Fabrication of poly(α-hydroxy acid) foam scaffolds using multiple solvent systems. J Biomed Mater Res 59(3):563–572 14. Maspero FA, Ruffieux K, Müller B, Wintermantel E (2002) Resorbable defect analog PLGA scaffolds using CO2 as solvent: structural characterization. J Biomed Mater Res 62(1):89–98 15. Niu W, Bai C, Qiu G, Wang Q (2009) Processing and properties of porous titanium using space holder technique. Mater Sci Eng A 506(1–2):148–151 16. Aydo˘gmu¸s T, Bor S¸ (2009) Processing of porous TiNi alloys using magnesium as space holder. J Alloy Compd 478(1–2):705–710 17. Mansourighasri A, Amirhossein N, Sulong AB (2012) Processing titanium foams using tapioca starch as a space holder. J Mater Process Technol 212(1):83–89

Processing of Haynes® 282® Alloy by Laser Powder Bed Fusion Technology Robert Otto, Vegard Brøtan, Amin S. Azar and Olav Åsebø

Abstract Haynes® 282® is a superalloy that offers excellent creep resistance up to 0.7 of its melting point. Promoted by aerospace industries, recommending a sound process window for additive manufacturing of this material is critical. In this study, the processing conditions were optimized by fabricating 84 test cubes using 36 different parameter sets. Thereafter, selected samples were solution annealed followed by a two-step aging for all. The cubes were optically analyzed for their porosity and defects to select the most promising parameters. Scanning electron microscopy (SEM) technique was employed to characterize those features. It was observed that selection of improper parameter combination has adverse effect on the porosity of the material and may cause systematic material cracking. Existing cracks were analyzed regarding their size and orientation pattern among selected and representative samples. Keywords Powder bed fusion · Additive manufacturing · Haynes 282 · Superalloy

R. Otto (B) · V. Brøtan · O. Åsebø SINTEF Manufacturing, S.P. Andersens veg 5, 7031 Trondheim, Norway e-mail: [email protected] V. Brøtan e-mail: [email protected] O. Åsebø e-mail: [email protected] A. S. Azar SINTEF Industry, Forskningsveien 1, 0373 Oslo, Norway e-mail: [email protected] © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_47

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Introduction Laser Powder Bed Fusion Laser Powder bed Fusion (PBF-LB) is a technology within the group of Additive Manufacturing (AM) processes. AM describes the fabrication of parts, as the name says, by adding material, often layer by layer [3]. In contrast to production by adding material is subtractive manufacturing, which includes processes like milling, drilling, turning and so on, in which the final shape of a part is achieved by removing material. Additionally, a combination of those conventional manufacturing methods with prior Additive Manufacturing is often called hybrid manufacturing. AM is divided into seven categories. Powder bed Fusion (PBF-LB) is such a category, with LPBF being a specific subcategory. This category describes a process where a layer of powder is spread on top of a substrate, thereby selectively being melted into a solid. Such melted stacks will grow into a shape that was described by a CAD model. Building a part layer by layer has the advantage of turning difficult three-dimensional shapes into simple 2D layers which are easier to handle and produce.

Nickle-Based Superalloys Out of the three families of superalloys, nickel-based superalloys are used for the most stringent high-temperature applications. For those, solid solution strengthening and precipitation hardening are the most common strengthening mechanisms, whereby precipitation hardening enables the important precipitates of gamma-prime (γ ) and gamma-double-prime (γ ). Due to their very small lattice misfit of around 0.1% with the austenitic nickel matrix (γ) and the resulting low surface energy, the face-centered cubic (fcc) gamma-prime became one of the most important precipitates for enabling high-temperature stability [2]. Both the heat treatment and chemical composition of an alloy play a crucial role in influencing the precipitation, whereby Al and Ti are known as γ forming elements [4, 5, 9]. As the content of these two components increases, more γ will be formed during aging and therefore lead to a higher strength. Nevertheless, with an increasing percentage of γ precipitates, problems during the fabrication can occur. Precipitation-hardened superalloys are generally susceptible to cracking during or after solidification (hot cracking and strain cracking, respectively) and as the content of precipitates raises, the forging range becomes narrower, requiring alternative fabrication processes like casting and powder metallurgy [2].

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Haynes 282 The nickel-based alloy Haynes 282 was developed in 2005. The wrought alloy was specifically designed for the best balance of strength and fabricability, applying γ strengthening [9]. Although the weldability and castability of Haynes 282 have been analyzed in several papers [6–8], there is currently no information about its fabrication with Laser Powder Bed Fusion. On the contrary, Inconel 718 is a prominent representative of already additively manufactured superalloys and is often described as one of the standard materials for high-end AM machines. In comparison to Haynes 282, it is mainly precipitate strengthened by gamma-double-prime (γ ). This leads to superior properties at lower temperatures on the one hand, yet quicker weakening at high temperatures on the other hand. Table 1 displays the difference between the chemical compositions of the two materials. Haynes 282 have ten times more cobalt and about triple the amount molybdenum, while Inconel balances around 15–20 times as much iron content and a niobium content of about five weight percentage. The range of standard materials for Powder bed Fusion is limited, although there are many metals that are under development or specially developed by powder manufacturers. Haynes 282 is a potential candidate to represent a valuable addition to the growing list of materials. The purpose of this study is to evaluate the manufacturing parameters of Haynes 282, as representative for new materials for PBF-LB. Following this approach, samples have been produced using different parameter sets and subsequently analyzed on selected properties. As will be explained, the porosity and microstructure of the fabricated parts are crucial to define the general setup for further fabrication. This paper will focus on the first analysis of those properties.

Table 1 Comparison of Haynes 282 and Inconel 718 composition, wt% Ni H282

57a

IN718 50− 55 a As

balance

* Maximum

B Balance

Cr

Co

Mo

Ti

Al

Fe

Mn

Si

C

B

1.5

1.5*

0.3

1.5*

0.06

0.005 –

B

0.35* 0.35* 0.08* 0.006 4.75– 5.5

20

10

8.5

2.1

17– 21

1*

2.8– 3.3

0.65– 0.2– 1.15 0.8

Nb

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Table 2 Haynes 282 composition of the powder compared to nominal values Ni

Cr

Co

Mo

Ti

Al

Fe

Mn

Si

C

B

Nominal

57a

20

10

8.5

2.1

1.5

1.5*

0.3

1.5*

0.06

0.005

Actual

56.94 19.48 10.14 8.49

2.11

1.56

1

0.15

0.08

0.06

0.004

a As

balance ∗ Maximum

Experimental Details Powder Material The feedstock used for this study was supplied by Haynes International and was sized at 11–45 μm. The nominal and actual composition of the Haynes 282 powder is given in Table 2. The nominal data in Table 2 is given by Praxair Surface Technologies, the producer of the powder. It was shown that the actual composition fulfilled the nominal requirements. All the samples in the study have been produced with all virgin powders.

Sample Design and Orientation The coordinate system used in this work is according to the definitions in ISO/ASTM 52900:2015 [3]. Hence, the X-direction is positive in the right-hand side of the machine. The coater is located in positive X, moving in negative X while spreading new layers of powder. The build platform consisted of a round Inconel 718 plate inserted into a larger square steel plate with silicone gaskets. All samples had 5 mm support structure to minimize the contamination between the Haynes 282 powder and the Inconel 718 build plate. The support structure had 0.5 mm squared hatching lines in X- and Y-directions with triangle formed point on the top to hold the part in place. Cubic samples for both builds were designed as 5 × 5 × 5 mm cubic samples placed on the described support structure. All samples were numbered in correlation to their manufacturing parameter set and placed in 5 mm distance from each other. This is shown in Fig. 1.

Manufacturing Parameters and Post-processing All samples were fabricated in a Concept Laser M2 using a 200 W Yb:YAG fiber laser. 99.9% pure nitrogen atmosphere was used as shielding gas during the fabrication to

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Fig. 1 The orientation of the cubes on the build plate

Fig. 2 Alternating scan strategy used in the manufacturing of test samples

avoid problems like oxidation. The influence of higher quality gasses, like argon, has not been proven apparent on the mechanical behavior [1]. The standard scan strategy in the machine is an island strategy with 5 × 5 mm islands. The samples were quite small, so the scan strategy was changed to an alternating scan like displayed in Fig. 2. In this strategy, the scanning direction is turned 90° every layer, always staying 45° to the X- and Y-axes. Two builds were conducted, whereby the manufacturing parameter for both is given in Table 3. Based on the powder analysis prior to the fabrication, process parameter for the first build was centered around the Concept Laser standard parameter set for Inconel 718. The parameter of the second build is based on the sample analysis of the first build, particularly on the results of porosity analysis. The chosen representative characteristic value to distinguish the used parameter sets was the volumetric energy density, calculated using the following: Evol  P/(v · t · h).

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The samples including support structure were heat treated after their removal from the build plate according to Haynes 282 data sheet. There is no approach of developing a heat treatment for PBF-LB processed Haynes 282 in this study.

Results Porosity Cubes for porosity tests have been prepared by electrical wire cutting, embedding and polishing. All parameter sets for both builds have been analyzed for porosity in their vertical and horizontal layer. The software ImageJ was applied with the Plugin called Otsu thresholding. As can be seen in Fig. 3, the porosity analysis of the first build displays a trend of increased densities at lower E vol , therefore, it was concluded that higher laser speeds were needed to find the near optimal value. It is clear that the density on the evaluated horizontal and vertical layers deviates from each other with the latter approaching lower relative density values. Additionally, the results of the vertical layers display a more volatile behavior. Nevertheless, homogeneous densities up to 99.7% could be achieved. The numbers presented a clear peak for the standard build parameter for Inconel 718 (E vol of 114.29 J/mm3 ), as both layer directions had the same density of around 99.5%. Following the results from the first build, the parameter range of build 2 was extended to a lower volumetric energy density, (E vol ). Due to the density peak around the Inconel 718 standard parameter set, a close set of parameters was also included in the second build. Figure 3 displays that it is still difficult to determine a clear density peak. However, the results demonstrate higher density in general for the second build, up to 99.9% homogeneous density. Once again, the Inconel 718 standard parameter set present a high homogeneous density for the vertical and horizontal layers of around 99.9% as well. Generally, the vertical layer exhibits a volatile behavior, although there are some flat peaks around 92 and 94.5 J/mm3 .

Table 3 Production parameters for build 1 and build 2 P (W)

v (mm/s)

t (mm)

h (mm)

Evol (J/mm3 )

Build 1

180

i[i + 1]  [0, …, 20], v  400 + 10i

0.03

0.105

95.24–142.86

Build 2

180

i[i + 1]  [0, …, 20], v  545 + 5i and i[i + 1]  [0, …, 2], v  495 + 5i

0.03

0.105

88.59–104.85 and 113.15-115.44

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Fig. 3 Comparison of density for build 1 and build 2

Crack Formation Figure 4 is representative for the best parameters used in this study. It displays the existence of several cracks, lying mainly perpendicular to the laser scanning direction. However, the cracks are repeating vertically along the build direction. The horizontal distance between the crack arrays is approximately 700 μm, while the cracks’ lengths are roughly 80–100 μm long. The crack length correlates with the horizontal length of grains in the XZ-plane, as can be seen in Fig. 4(2). Most cracks are equally shaped, therefore likely to be caused by the same reason. Also, cracks seem to start in between two overlapping areas of grains and not, as might expected, inside the overlapping areas in the XZ-plane. The size of grains correlates directly with the distance of overlapping areas, which will result in a mosaic pattern in the XY-plane.

Conclusion and Further Work It has been proven that the PBF-LB processing of Haynes 282 is principally feasible. Samples with up to 99.9% of homogeneous density have been produced. These high densities appear in several production parameter sets, which is positive for the reliability of the production. The concluded best choice from this study is the 94.5 J/mm3 parameter. Nevertheless, most analyzed samples have shown signs of crack formations. The formations consist of several horizontal cracks ordered above each other in the direction of the build direction. Based on our early speculations,

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Fig. 4 SEM images of the XZ-plane in a mid-cut of the part

the appearance of the evolved cracks is related to the laser scan strategy and possibly caused by hot cracking phenomena. Further studies on the crack development will be conducted to explain their creation mechanisms and influence. This will include SEM, EBSD and simulation approaches. Additionally, the mechanical properties will be evaluated in the future. Acknowledgements The work presented in this paper is a part of the Norwegian national research project “Material Knowledge for Robust Additive Manufacturing”, which is supported by GKN Aerospace Norway, Nammo Raufoss, Kongsberg Automotive, Sandvik Teeness and OM BE Plast. The project is funded by The Research Council of Norway (Grant no. 248243).

References 1. Amato KN, Gaytan SM, Murr LE, Martinez E, Shindo PW, Hernandez J, Collins S, Medina F (2012) Microstructures and mechanical behavior of Inconel 718 fabricated by selective laser melting. Acta Mater 60(5):2229–2239 2. Campbell F (2008) Elements of metallurgy and engineering alloys. ISBN: 9780871708670 3. ISO (2015) ISO/ASTM 52900:2015 Additive manufacturing—part 1: terminology 2015(X) 4. Kim HT, Chun SS, Yao XX, Fang Y, Choi J (1997) Gamma prime (γ ) precipitating and ageing behaviors in two newly developed nickel-base superalloys. J Mate Rials Sci 32:4917–4923 5. Maniar GN, Bridge JE, James HM, Heydt GB (1970) Correlation of gamma-gamma prime mismatch and strengthening in Ni/Fe-Ni base alloys containing aluminum and titanium as hardeners. Metall Trans 1(1):31–42 6. Osoba LO (2012) A study on laser weldability improvement of newly developed Haynes 282 superalloy 7. Osoba LO, Ding RG, Ojo OA (2012) Microstructural analysis of laser weld fusion zone in Haynes 282 superalloy. Mater Charact 65:93–99 8. Osoba LO, Ojo Oa (2012) Influence of laser welding heat input on HAZ cracking in newly developed Haynes 282 superalloy. Mater Sci Technol 28(4):431–436 9. Pike LM, International H, Ave WP (2008) Development of a fabricable gamma prime (γ ) strengthened superalloy. In: Superalloys 2008, pp 191–200

Part IX

Advanced High-Strength Steels III

Tensile Deformation Behavior of 1 GPa-Grade TRIP-Aided Multi-microstructure Steels Studied by In Situ Neutron Diffraction Noriyuki Tsuchida, Takaaki Tanaka and Yuki Toji

Abstract Tensile deformation behavior and the TRIP effect of 1 GPa-grade TRIPaided multi-microstructure (TRIP) steels were studied by in situ neutron diffraction experiments during tensile test. The effect of retained austenite (γ R ) shape on TRIP effect in the 1 GPa-grade TRIP steels was focused on. In the static tensile tests, the 1 GPa-grade TRIP steel with the γ R shape of needle-like showed better uniform elongation, whereas the tensile strength was almost the same as the TRIP steel with the γ R shape of blocky. The reasons for the better uniform elongation of the needle-like TRIP steel were discussed from the viewpoints of phase strain and deformation-induced martensitic transformation behavior obtained by the neutron diffraction experiments. As a result, the difference of phase strains and the transformation rate are found to play an important role in the better uniform elongation. Keywords TRIP · Uniform elongation · Neutron diffraction

Introduction Transformation-induced plasticity (TRIP) is a strengthening mechanism that can obtain better elongation and high strength due to the deformation-induced martensitic transformation of retained austenite (γ R ) [1, 2]. The TRIP effect is exploited for the manufacture of various steels such as TRIP-aided multi-microstructure (TRIP) steel and one potential application of TRIP steel is steel sheets for automobiles. N. Tsuchida (B) University of Hyogo, Himeji, Japan e-mail: [email protected] T. Tanaka · Y. Toji JFE Steel, Chiba, Japan e-mail: [email protected] Y. Toji e-mail: [email protected] © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_48

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We studied tensile deformation behavior with a wide range of strain rate of 1 GPa-grade TRIP steels with different γ R shapes obtained by a 0.3C steel [3]. As a result, the TRIP steel with γ R shape of needle-like showed the tensile strength of 1 GPa and the total elongation of 40% in the static tensile test at 296 K [3]. In this time, when the mechanical properties were compared between the two TRIP steels with different γ R shapes, the TRIP steel with γ R shape of needle-like showed better uniform elongation whereas their tensile strength were approximately 1 GPa independent of the γ R shape. In this study, in order to discuss reasons for the better uniform elongation of the TRIP steel with γ R shape of needle-like, in situ neutron diffraction experiments during tensile test at room temperature were conducted.

Experimental Procedures In this study, two types of TRIP steel with different γ R morphologies were prepared from 0.3C-1.5Si-2Mn (mass%) steel and named needle-like γ R steel and blocky γ R steel [3]. The matrix of the needle-like γ R steel in is martensite and partially bainite, while that of the blocky γ R steel is ferrite and bainite. The volume fraction of γ R was 24.6% for the needle-like γ R steel and 22.9% for the blocky γ R steel. In addition, the volume fraction of ferrite and bainite according to the point counting method in the optical images of the blocky γ R steel was 39.7% and 37.4%, respectively. From the TRIP steels, plate tensile test specimens with the gage length of 55 mm, width of 6 mm and thickness of 1.4 mm were prepared. Here, the tensile loading direction was adjusted being parallel to the rolling direction. The in situ neutron diffraction experiments during tensile test were conducted at TAKUMI, a high resolution and high-intensity time of flight neutron diffractometer for engineering sciences at Materials and Life Science Experimental Facility of Japan Proton Accelerator Research Complex (J-PARC) [4]. The specimen was mounted horizontally in a loading machine installed at TAKUMI, in such a way that the neutron diffraction patterns in the axial and transversal directions were measured simultaneously using two detector banks. The detailed information about TAKUMI is described in Ref. [4]. The tensile deformation was conducted with an initial strain rate of 6.1 × 10−6 s−1 in elastic regime, and with an initial strain rate of 1.8 × 10−5 s−1 in plastic regime. Diffraction patterns were then extracted periodically with 300 s accumulating time for different applied strain or stress values. Data analyses were performed by a single peak fitting method and a multi peak fitting method using a Rietveld software, which is so-called Z-Rietveld [4].

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Results and Discussion Figure 1 shows nominal stress–strain curves of the 1GPa-grade TRIP steels obtained by the static tensile tests at 296 K [3]. Tensile strength of the present two TRIP steels are almost the same of about 1 GPa, but the 0.2% proof strength, uniform elongation and total elongation of the needle-like γ R steel are better than those of the blocky γ R steel. Especially, it is interesting that the uniform elongation of the needle-like γ R steel is about 5% larger than the blocky one though the tensile strength of the both TRIP steel showed approximately 1 GPa. Figure 2 shows true stress and work-hardening rate as functions of true strain in the 1GPa-grade TRIP steels. The work-hardening rate of the blocky γ R steel was larger than that of the needle-like

1200

Needle-like

R

1000

Nominal stress (MPa)

Fig. 1 Nominal stress–strain curves of the 1 GPa-grade TRIP steels obtained by static tensile test at 296 K

800

Blocky

R

600

400

200

Strain rate = 3.3x10-4 s-1 0

0

0.05

0.1

0.15

0.2

0.25

0.3

0.35

0.4

Fig. 2 True stress and work-hardening rate as functions of true strain of the 1 GPa-grade TRIP steels obtained by static tensile test at 296 K

True stress or work-hardening rate (MPa)

Nominal strain 4000 3500

Blocky γR

3000 2500

Needle-like γR

2000

dσ/dε

1500 1000

σ

500 0

0

0.05

0.1

0.15

True strain

0.2

0.25

0.3

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γ R steel at small strain range until about 0.15. At the true strains more than 0.15, the work-hardening rate of the needle-like γ R steel became larger. The work-hardening rate of the needle-like γ R steel remains almost the same when true strain is between 0.1 and 0.2. Figure 3 shows phase strains of ferrite (α), austenite (γ ) and martensite (α  ) phases as functions of true stress in the needle-like and blocky γ R steels. Here, YS means the yield strength (0.2% proof strength) of each TRIP steel. The phase strain of α  phase is the largest in the both TRIP steel and the change of phase strain for α  was almost independent of the γ R shape. On the other hand, in terms of changes for α and γ phases, the phase strains between α and γ were different before YS and the difference at YS for the blocky γ R steel was larger than the needle-like γ R steel. The difference of phase strains between α and γ for the blocky γ R steel was larger than the needle-like γ R steel until the true stress of 1,100 MPa. But that for the needle-like γ R steel became larger at the true stress more than 1,100 MPa. On the other hand, the phase strain for γ in the needle-like γ R steel became larger than the blocky γ R one when true stress was more than 800 MPa and the increase of phase strain for γ was also larger than the blocky γ R steel. The difference of phase strains is associated with work-hardening behavior [5, 6]. The larger difference of phase strains, the larger the work-hardening becomes. And such the difference of phase strains was also associated with the changes of work-hardening rate as seen in Fig. 2. Figure 4 shows the volume fraction of deformation-induced α  as a function of true strain (a) and true stress (b). Arrows in Fig. 4b are the yield strength (YS) of each TRIP steel. In Fig. 4a, the volume fraction of deformation-induced α  at a given 0.015

Solid: Needle-like γR Opened: Blocky γR

YS (Needle-like) YS (Blocky)

0.01

Phase strain

Martensite Austenite

0.005

Ferrite

0

0

200

400

600

800

True stress (MPa)

1000

1200

1400

Fig. 3 Phase stains of ferrite (α), austenite (γ ) and martensite (α  ) as functions of true stress obtained by the in situ neutron diffraction experiments during tensile deformation

Volume fraction of deformation-induced martensite

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0.25

(a)

(b)

Yield strength

0.2

0.15

Blocky

Needle-like R

R

0.1

Blocky Needle-like

0.05

0 0

0.05

0.1

0.15

True strain

0.2

0.25

R

R

0 0.3

500

1000

1500

True stress (MPa)

Fig. 4 Volume fraction of deformation-induced martensite as a function of a true strain and b true stress in the 1 GPa-grade TRIP steels

true strain of the blocky γ R steel was larger when compared to the needle-like γ R steel. But the volume fraction of deformation-induced α  at the maximum load point was almost the same as about 10%. On the other hand, in Fig. 4b, the deformationinduced transformation started after YS in the needle-like γ R steel whereas that in the blocky γ R steel occurred before YS. This means the deformation stability of γ R in the needle-like γ R steel is higher than the blocky one. Furthermore, the volume fraction of deformation-induced α  increased linearly according to true stress and the slope of the needle-like γ R steel was larger than that of the blocky one. That is, the transformation rate with true stress of the needle-like γ R steel is larger compared to the blocky one. Here, the transformation rates of the needle-like and blocky γ R steels are 0.0179 %/MPa and 0.0125 %/MPa, respectively. TRIP effect is effective in the case that the γ R transforms to deformation-induced martensite actively in the latter part of deformation from the past studies [2, 3]. This means the volume fraction of γ R and the transformation rate at the large true stress (or true strain) range are important factors for the better uniform elongation of the needle-like γ R steel.

Summary 1. The phase strain of γ for the needle-like γ R steel continued increasing until about 1,000 MPa and the difference of phase strains between α and γ also became larger at the large true stresses. 2. The volume fraction of deformation-induced martensite can be summarized as a function of true stress in both TRIP steels. The needle-like γ R steel shows a higher mechanical stability of γ R than the blocky γ R steel and the transformation rate was also larger.

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3. The reasons for the larger uniform elongation of the 1 GPa-grade TRIP steel with γ R shape of needle-like are the deformation-induced transformation behavior and the deformation behavior of γ and α phases at larger stresses or strains. Especially, the stress of γ phase and its change during tensile deformation play important roles in the both of stress–strain relationship and the deformationinduced transformation behavior of γ R .

References 1. Han HN, Oh C-S, Kim G, Kwon O (2009) Design method for TRIP-aided multiphase steel based on a microstructure-based modelling for transformation-induced plasticity and mechanically induced martensitic transformation. Mater Sci Eng A 499:462–468 2. Angel T (1954) Formation of martensite in austenitic stainless steels: effects of deformation, temperature, and composition. J Iron Steel Inst 177:165–174 3. Tsuchida N, Okura S, Tanaka T, Toji Y (2017) High-speed tensile deformation behavior of 1 GPa-grade TRIP-aided multi-phase steels 58(5):978–986 4. Harjo S, Tsuchida N, Abe J and Gong W (2017) Martensite phase stress and the strengthening mechanism in TRIP steel by neutron diffraction. Sci Rep 7(15149). https://doi.org/10.1038/ s41598-017-15252-5 5. Tomota Y, Tokuda H, Adachi Y, Wakita M, Minakawa N, Moriai A, Morii Y (2004) Tensile behavior of TRIP-aided multi-phase steels studied by in situ neutron diffraction. Acta Mater 52:5737–5745 6. Ojima M, Adachi Y, Tomota Y, Ikeda K, Kamiyama T, Katada Y (2009) Work hardening mechanism in high nitrogen austenitic steel studied by in situ neutron diffraction and in situ electron backscattering diffraction. Mater Sci Eng A 527:16–24

Development of Advanced High-Strength Steels for Automobile Applications Francys Barrado, Tihe Zhou, David Overby, Peter Badgley, Chris Martin-Root, Sarah Zhang and Rich Zhang

Abstract Stelco has developed a suite of Advanced High-Strength Steels (AHSS) grades with tensile strength greater than 1000 MPa to meet standard automotive specifications and for unique customer requirements. These grades were optimized by correlating chemical composition and processing parameters with microstructures and mechanical properties. Dual-Phase 980 (Stelco trademarked STELMAXTM 980DP), Multiphase 1180 (STELMAXTM 1180MP), Martensitic 1300 (STELMAXTM 1300M) and 1500 (STELMAXTM 1500M) products met strength and formability requirements with excellent flatness and consistent mechanical properties across the entire strip length and width by using hydrogen quench continuous annealing technology. Keywords Advanced high-strength steels · Microstructure · Mechanical properties · Applications

Introduction Advanced High-Strength Steels (AHSS) have been applied extensively to the bodyin-white starting in the early 2000s. There are several classes of AHSS: dual-phase (DP), transformation-induced-plasticity (TRIP), multiphase phase (MP), twinnedinduced-plasticity (TWIP) and martensitic steel (MS). The utilization of these steel grades is driven by more stringent passenger safety and fuel consumption requirements. The new Corporate Average Fuel Economy (CAFE) 2025 was set at 54.5 mpg by the Obama government, although this number is under review by Environmental Protection Agency (EPA) [1, 2]. These steel grades can be manufactured by different processes such as hot-dip galvanizing, batch annealing and continuous annealing with different cooling modes such as rapid gas jet cooling and water quench. This paper will discuss recent developments of different AHSS grades produced F. Barrado · T. Zhou (B) · D. Overby · P. Badgley · C. Martin-Root · S. Zhang · R. Zhang Research Department, Stelco Inc., 386 Wilcox Street, Hamilton, ON L8L 8K5, Canada e-mail: [email protected] © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_49

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on a continuous annealing line with hydrogen quench technology. Grades developed were Dual-Phase 980 (Stelco trademarked STELMAXTM 980DP), Multiphase 1180 (STELMAXTM 1180MP), Martensitic 1300 (STELMAXTM 1300M) and 1500 (STELMAXTM 1500M). A correlation between chemical composition and processing parameters from the hot strip mill, through to the cold mill and continuous annealing line will be presented.

Chemistry Design The chemistry design of AHSS is essential to meet the required strength levels while balancing formability, weldability and edge stretchability, as well as the annealing profile, i.e. critical hydrogen quench cooling rate, annealing time and temperature at maximum line speed for each given cross section has to be taken in consideration. A comparison of chemistry design for different AHSS grades developed at Stelco Inc. is listed in Table 1. Control of phase transformation is critical to the development. It is necessary to promote the coexistence of different microstructural constituents; their individual mechanical behavior and their mutual interaction develop the desired strength–ductility balance. Proper control of those multiphase microstructures is accomplished through chemistry design and processing control. DP steel is composed of soft ferrite matrix and 10–70% of hard martensite islands depending on DP grades. To achieve the unique combination of mechanical properties for DP steel, carbon is restricted from 0.10 to 0.20%; 0.3–0.7% Si, 1.0–2.0% Mn. Up to 0.4% Cr and/or Mo are added to stabilize austenite and strengthen the ferrite. In some cases, Nb, Ti and V are used to refine the microstructure and promote ferrite formation [3]. The presence of soft phase and hard phase in the dual-phase microstructure has high sensitivity to edge fracture i.e. low edge stretchability. To overcome the low edge stretchability, STELMAXTM 1180MP is designed to reduce the differences between soft and hard constituents. Thus, Cr and Mo are added to suppress ferrite formation. Martensitic steels are the hardest type of AHSS which have high tensile strength and high yield-to-tensile strength ratio. The strength of STELMAXTM 1300M and 1500M is controlled mostly by carbon content; however, other alloying elements such as Mn, Cr, Si and Mo could be also added to achieve the required strength, ductility, bendability and weldability [4, 5].

Table 1 Chemical composition ranges (%) Grade

C

Mn

P

Si

S

Cr + Mo

Ti + Nb

STELMAXTM 980DP

0.10–0.20

1.50–2.00

Fr . μ  · ln (1) θ Fr

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Evaluation of Experimental Results The results of the presented investigations refer exclusively to experiments with liquid CO2 as a volatile lubricant. The CO2 was taken from a gas cylinder with a riser pipe so that liquid CO2 could flow into the radius insert at approx. 60 bar (6 MPa, at a temperature of 20 °C). In this case, the phase change to the gaseous state takes place shortly before leaving the microholes [6]. A summary of the experimental parameters and settings can be found in Table 1.

Influence of Retention Force on the Coefficient of Friction In order to set different surface pressures at the radius insert, the retention force was adjusted within four levels (cf. Table 1). In the investigation presented in this paper, the injection angle was kept constant at 45°. All parameter combinations of executed friction experiments were repeated five times in order to statistically evaluate the results. The repeatability of these experiments showed very small fluctuations in the measured values of less than 2% across all parameter combinations. This can also be seen in Fig. 4 from the low standard deviation values of the error bars. As expected, the friction coefficients at the tool radius insert are relatively high compared to the flat strip drawing tests (see paragraph 4). This corresponds to the real forming process, where the high surface pressure distributions at tool radii also result in a significant increase of the friction coefficients. As can be seen in Fig. 4, a mean friction coefficient of μ  0.26 could be determined at a retention force of 3.4 kN. With a retention force of 6.1 kN, a friction value of μ  0.19 was measured on average. Here, the clear tendency of the friction

Table 1 Experimental parameters and settings for deflected strip drawing tests

Sheet metal material

DC05+ZE

Tool material

1.2379 fully hardened to HRC60, polished

Tool radius insert

R5

Volatile media

CO2

Gas pressure

60 bar  6 MPa

Microholes

11 microholes single-line arranged, diffuser

Injection angle

45°

Drawing velocity

100 mm/s

Sheet metal strip

590 mm × 60 mm

Retention force Fr

3.4, 4.0, 5.0 and 6.1 kN

Injection angle α

30°, 35°, 40°, 45°, 50°, 55°, and 60°

G. Reichardt and M. Liewald

coefficient of friction μ [-]

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0.27

Coefficient of friction at radius 5 mm using CO2

0.25 0.23 0.21 0.19 0.17 0.15 3.00

3.50

4.00

4.50

5.00

5.50

6.00

6.50

retention force Fr [kN]

Fig. 4 Coefficient of friction at radius insert 5 mm using CO2 as lubricant substitute, injection angle 45° Fig. 5 Definition of injection angle α of volatile fluids

Drawing direction

Sheet metal strip

α Radius insert

coefficients is interesting, as friction decreases in this range of values with increasing surface pressure. As known from the literature, the value of the coefficient of friction diminishes with increasing surface pressures and is also strongly influenced by the dimensions of the radius inserts [7].

Optimization of Media Injection Angle In order to characterize the influence of the injection angle on the friction behaviour of the tribological system at the most heavily loaded areas of a deep-drawing tool, i. e. at the tool radii, a separate investigation was carried out. The aim was to optimize the injection angle in order to achieve the lowest possible coefficients of friction at different restraining forces. The definition of the injection angle α is shown in Fig. 5 and was changed for this investigation between 30° and 60° in steps of 5°. The results of this investigation are shown in Fig. 6. The aim of this procedure was to define a value range of the injection angle for which the lowest possible coefficients of friction can be achieved for different restraining forces. In that way, an optimal value for the injection angle could be found independent of the restraining forces

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coefficient of friction μ [-]

Coefficient of friction at radius insert 5 mm with varied injection angles for different retention forces 0.35 0.33 0.31 0.29 0.27 0.25 0.23 0.21 0.19 0.17 0.15

Low friction area

30

35

40

45

50

55

60

injection angle [°] Fr 3,4kN

Fr 4kN

Fr 5kN

Fr 6kN

Fig. 6 Optimization of the coefficient of friction for different injection angles using CO2 at different retention forces

and thus also the corresponding surface pressures. As can be seen in Fig. 6, the area with low coefficients of friction is between 40° and 45° of the injection angle. The tendency shows that with increasing restraining forces the coefficients of friction decrease from a maximum of 0.31 to a minimum of 0.19. This also coincides with the results in Chapter 3.1 and with the findings in [7]. The results of the experiments with retention forces (Fr ) of 3.4 and 4 kN show the same tendency. With small injection angles, the coefficients of friction remain relatively constant. This means that the volatile medium introduced into the gap between tool radius and sheet metal strip can be supplied at an angle between 30° and 45° without any influence on friction. If, however, the injection takes place at angles greater than 45° and small restraining forces of 3.4 kN, a significant increase in friction can be seen, which is also noticeable by the occurrence of zinc adhesions on the radius inserts. The occurrence of these adhesions at larger injection angles also explains the relatively high dispersion of these values. For the same reason, the experiments at Ft equal to 4 kN were stopped at 55°. At a restraining force of 5 kN, an exactly reversed behaviour can be observed. Up to an injection angle of 35°, relatively high coefficients of friction occur and from 40° the values remain fairly constant at approx. 0.22. It can, therefore, be assumed that with higher surface pressures a better sealing (gaseous lubricant) of the friction zone occurs, which has a positive effect on friction above an injection angle of approx. 40°. This behaviour is clearly evident with a retention force of 6 kN. Here, a constantly relatively low coefficient of friction occurs over the entire angular range. In order to further validate the influence of the zinc coating, which may also have friction-reducing effects, on these results, strip drawing experiments with uncoated DC05 strips will be carried out in further investigations. In addition, the arrangement of the microholes will be supplemented with a second row of holes in order to improve friction in the area of small injection angles, especially at lower retention forces.

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Comparison of Flat and Deflected Strip Drawing Experiments/Discussion of Deviations In order to improve the comparability of the results from the strip drawing experiments with and without deflection, the surface pressure occurring is usually used as a reference value. In flat strip drawing experiments, the surface pressure is constantly distributed over the friction surface, which greatly simplifies the determination of the surface pressures. In the strip drawing experiments with deflection, however, no constant surface pressure occurs at the contact surface between the sheet metal strip and the radius insert. Here the surface pressure distribution is similar to the analytical calculation of plain bearings according to Hertz (Hertz surface pressure), although the condition of a purely elastic behaviour of both contact partners is not given. However, since a clear plastification occurs in the sheet metal strip, Hertz’s approach cannot be applied here. For a first and simple comparison, a very simple method of calculating the surface pressure at the radius was used.

Analytical Determination of the Averaged, Projected Surface Normal Pressure As shown in Fig. 7, the averaged, projected surface pressure (pnormal ) is calculated using the projected surface Aproj and the corresponding orthogonal force components of Ft and Fr . The value determined in that way represents an average surface pressure along the contact surface. In reality, the surface pressure distribution at the contact start and contact end are zero and at the center of the radius at the highest. Furthermore, the maximum value in the middle is clearly above the average value and is limited by the plastification of the sheet material. This system behaviour also

Sheet metal strip

Drawing direction

Ft Aproj

Ft : tractive force Fr : retention force Aproj: projected contact surface of radius insert

Radius insert Fres Fr

Fres = Ft cos(45°) + Fr cos(45°) Fres Aproj

Fig. 7 Calculation of the averaged, projected surface normal pressure at the radius insert

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Fig. 8 Main effects on the coefficient of friction μ determined with flat strip drawing investigations at 5 MPa (lower, blue values) and 15 MPa (higher, green values) surface pressure, drawing speed 100 mm/s, lubrication medium CO2 [5] (Color figure online) Table 2 Averaged, projected surface normal pressure of deflected strip drawing experiments

Deflected strip drawing test, injection angle 45° Retention force Fr (kN)

Coefficient of friction μ (–)

pnormal (MPa)

3.44

0.260

14.4

4.00

0.242

16.4

5.00

0.213

20.0

6.13

0.191

24.1

explains the clear differences in friction coefficients between flat and deflected strip drawing experiments as can be seen in Fig. 8 and Table 2. For future investigations, it is planned to determine the occurring contact normal stress distributions by means of simulations. In this way, the maximum value of the contact normal stress in the middle of the die entry radius and the contact normal stress can be determined associated for incipient adhesions.

Comparison of the Results As presented, the friction ratios determined by means of flat strip drawing experiments [5] showed significantly lower coefficients of friction (cf. Fig. 8: pnormal 5 MPa: 0.022–0.051; pnormal 15 MPa: 0.065–0.074) than the investigations with the new strip drawing experiments with deflection.

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When comparing the friction values of flat and deflected strip drawing experiments, the values of the higher surface pressure of the flat strip drawing experiments (15 MPa, μ  0.065–0.074) show considerably lower friction values than the results of the deflected strip drawing experiments (14.4 MPa, μ  0.242). This significant difference can be explained by the difference in the actually occurring surface pressure at the radius inserts of the strip drawing experiments with deflection. The non-constant distribution of the contact normal stress over the radius cross-section results in significantly increased surface pressure values in the middle of the radius. The determination of the contact normal stress distribution over the radius cross section of these values by FE-simulation will be part of the following investigations.

Outlook The presented investigations show the feasibility of a tribological system with volatile media as lubricant substitute. Even with relatively high retention forces, a simple microhole arrangement can prevent wear due to adhesion or cohesion. However, the test results show that some areas of the new tribo-system still need to be investigated in more detail. The comparison of the friction values of flat and deflected strip drawing experiments clearly shows the different friction ratios of both testing variants. The main reason for this is the uneven normal stress distribution over the radius crosssection. Therefore, the knowledge of the contact normal stress distributions and thus the maximum surface pressures occurring at the radius inserts is extremely important for the various friction conditions and will be content of further investigations and comparisons. Furthermore, for a complete characterization of this new tribo-system further influencing factors have to be investigated. It is planned to vary the radius insert size, the arrangement and number of microholes, the drawing speed, the gas supply pressure and the gas type in order to determine their influence on the friction behaviour. With the planned investigations a more comprehensive understanding of the system behaviour of this new tribo-system can be achieved. Acknowledgements The scientific investigations of this paper are funded by the German Research Foundation (DFG) within the priority program SPP 1676 Dry Metal Forming—Sustainable Production by Dry Processing in Metal Forming. We thank the German Research Foundation (DFG) for the funding of this research project.

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References 1. Bay N (2011) Trends and visions in metal forming tribology. In: Special edition: 10th international conference on technology of plasticity, ICTP, pp 15–26 2. Vollertsen F, Schmidt F (2014) Dry metal forming: definition, chances and challenges. Int J Precis Eng Manuf Green Technol 1(1):59–62 3. Bay N et al (2010) Environmentally benign tribo-systems for metal forming. CIRP Ann Manuf Technol 59:760–780 4. Liewald M et al (2015) Tribosysteme für die Kaltumformung auf der Basis von flüchtigen Schmiermedien und laserstrukturierten Oberflächen. Dry Met Form OAJ FMT 1:22–33 5. Woerz C, Zahedi E, Umlauf G, Liewald M, Weber R, Graf T, Tovar GEM (2017) Tiefziehen eines U-Profils mit flüchtigen Medien als Schmierstoffersatz. Dry Met Form OAJ FMT 3:50–61 6. Umlauf G, Zahedi E, Woerz C, Barz J, Liewald M, Graf T, Tovar GEM (2016) Grundlagenuntersuchungen zur Herstellung von Lasermikrobohrungen in Stahl und dem Ausströmverhalten von CO2 als Trockenschmiermedium. Dry Met Form OAJ FMT 2:18–24 7. Papaioanu A (2016) Einsatz eines neuartigen Verfahrens zum kombinierten Recken und Tiefziehen von Außenhauptbeplankungen aus Feinblech, Beiträge zur Umformtechnik 81. Dissertation, University of Stuttgart, Institute for Metal Forming Technology

Investigation of Friction and Adhesion Behavior of Textured Workpieces and Coated Tools Under Dry Tribological Contact Rafael Hild, Robby Mannens, Daniel Trauth, Patrick Mattfeld, Thomas Bergs, Dennis C. Hoffmann, Nathan C. Kruppe, Tobias Brögelmann and Kirsten Bobzin Abstract Processes in cold forming are accompanied by high process loads, which are reduced by means of lubricants. Lubricants are ecologically and economically questionable and increasingly restricted by legislation. In a current research project, the substitution of lubricants by coatings for tools and texturing of workpieces in a full forward extrusion process is investigated. Exact friction values of combinations of texturing and coating are necessary to investigate the process numerically and experimentally. For solid forming, the ring compression test has proven to be a valid analogy test for determining the friction factor. In this paper, the latest investigations concerning the dry tribological contact of coated and uncoated tools with different textured workpiece surfaces are shown in the ring compression test. The flow and adhesion behavior of the active partners under loads of cold forming are investigated. Furthermore, the friction factor for numerical investigations of full forward extrusion process is determined. Keywords Cold forming · Tribology · Surface integrity

Introduction Cold forming comprises the most important mass production processes. Cold forming processes are characterized by high material utilization, high product quality and short process cycles [1]. At the same time, the material properties are influenced by work hardening. These process characteristics are ideal for the mass production of highly stressed components [2]. Especially, the full forward extrusion plays an important role. During full forward extrusion, high loads occur depending on the R. Hild (B) · R. Mannens · D. Trauth · P. Mattfeld · T. Bergs Laboratory for Machine Tools and Production Engineering (WZL), RWTH Aachen University, Aachen, Germany e-mail: [email protected] D. C. Hoffmann · N. C. Kruppe · T. Brögelmann · K. Bobzin Surface Engineering Institute (IOT), RWTH Aachen University, Aachen, Germany © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_151

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tapering ratio. Lubricants, surface structuring on workpieces and coatings on tools are used to ensure that the tools withstand the high process loads at high batch sizes [LANG08]. Above all, solid lubricants allow a reduction of process loads and thus protect the tool from damage or wear. Solid lubricants such as phosphating are combined with surface structuring by shot peening. Shot blasting enlarges the surface of the workpiece. A larger surface area allows a higher lubricant absorption capacity of the workpiece and thus continuously supplies the die surfaces with lubricant [3]. However, the use of lubricants is ecologically and economically questionable and has been increasingly restricted by legislation in recent years [4]. Due to this change, the research initiative SPP1676 is investigating lubricant-free forming. Reducing or even eliminating lubricants in the process would reduce process steps, which in turn would result in economic savings [5]. Furthermore, in addition to the economic savings, the ecological burdens based on the use of lubricants would also be reduced. However, the reduction of lubricants inevitably leads to an increase in process loads and could damage the tool and worsen the overall process. Dry forming is still largely unexplored in the field of solid forming [6]. In the field of sheet metal forming, various researches have been carried out in the direction of dry forming. Using DLC-coated deep-drawing dies, Murakawa achieved a dry sheet metal forming process [7]. Murakawa et al. expanded their findings by studying different tribological conditions and were able to show that siliconefree films further reduce friction [8]. Osakada et al. investigated the influence of the roughness of the tool surface in the dry tribological system and found that a rougher surface leads to greater friction [9]. Kataoka et al. investigated the influence of ceramic tools on the dry deep-drawing process and were thus able to achieve an equivalent process to the lubricated state for deep-drawing low-alloy steels [10]. Tamaoki et al. developed electrically conductive ceramic tools with which similar deep-drawing aspects as in the lubricated state could be achieved [11]. These findings of sheet metal forming cannot be transferred to cold forming due to the higher process loads and higher surface enlargement. Coatings on tools and the use of solid lubricants on previously blasted workpieces are used in cold forming to reduce friction [12]. Due to the omission of the lubricant, the mechanisms of action of the lubricant are also omitted. These must be substituted by the other active partners of the tribosystem. Up to now, the influence of coatings on tools or structures on workpiece surfaces under loads of cold forming on friction has not been investigated. A validated method for determining the friction coefficient of a tribological system for cold forming is the ring compression test (RCT) [13]. The ring upsetting test is an analogy test under laboratory conditions with specimens based on the upsetting method [14]. A circular annular specimen is compressed between two plane-parallel upsetting plates. The initial geometry is defined by outer diameter D0 to inner diameter d 0 to height h0 according to formula 1 as follows [15]: D0 : d0 : h 0  6 : 3 : 2.

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Sofuoglu et al. investigated the influence of the material properties and the test conditions on the flow conditions and the determined friction value with the aid of the ring compression test [16]. They found that each material has its own calibration curves for the ring compression test. Dehghan et al. investigated the applicability of the ring upsetting test for hot forging. They found that the ring upsetting test cannot be used under hot forging conditions [17]. Bay et al. investigated alternative ring geometries for evaluating friction under low contact normal stresses. They showed that the changed geometry achieves more precise friction values under low normal contact stress [18]. Mattfeld investigated the influence of different lubricants on the friction coefficient in the ring compression test. The aim was to find an environmentally friendly lubricant with similar or lower coefficients of friction than conventional lubricants. Spray oiling by MoS2 showed the lowest coefficient of friction [19]. The latest approach uses self-lubricating coatings on tools and surface textures on workpieces in order to enable a first dry metal forming process, see Fig. 1 [20]. In former investigation, the influence of surface textured workpieces on the use of self-lubricating PVD-coated tools is done [21]. The boundary conditions of the tribosystem are either changed by the use of a self-lubricating coating (Cr,Al)N + X:S (X  Mo, W) on tools or surface textures by means of shot peening. Tribosystem for dry metal forming smooth Tool F

rough PVD-coating (Cr,Al)N+X:S (X=Mo,w)

Surface texture

textured v Workpiece

Fig. 1 Research objectives for a novel tribosystem for dry metal forming in view of the tool from the surface engineering view (IOT) and surface texturing view (WZL) [20]

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The aim of this work is to investigate the influence of different surface structures and the influence of a tool coating on the friction coefficient and the resulting flow conditions. Thus, the optimal combination with regard to friction reduction is determined and the understanding of the mechanisms of action between structured semi-finished product surfaces and coated tools is improved.

Previous Work The current research is realized in the second phase of the priority program SPP 1676 of the German Research Foundation. In phase I, new approaches to enable dry metal forming processes were investigated initially using stationary loads. In this section, an explanation of a novel pin-on-cylinder tribometer, which was used to analyze the influence of surface textures on the frictional shear stress, is given. To do so, a long frictional path has to be analyzed in order to ensure that an untreated surface texture is in contact with the pin along the frictional path. The experimental results are further complemented by the development of a numerical FE model to analyze the influence of different surface textures on the friction shear stress. These results are correlated with shot peening parameters of the peening process to distinguish a parameter set to reduce the frictional shear stress the most for each shot peening medium. After the first step of tribometer tests, the surface textures are applied on specimens for an extrusion process in order to distinguish the most suitable surface texture to reduce friction and by that the punch force.

Deposition Process of the (Cr,Al)N + Mo:S Coating The (Cr,Al)N + Mo:S-coatings were deposited by means of the hybrid direct current magnetron sputtering/high-power pulsed magnetron sputtering (dcMS/HPPMS) technology in an industrial scale coating unit CC800/9 Custom, CemeCon AG, Wuerselen, Germany. The unit is constructed with a chamber volume of V  1 m3 . For the deposition, five CrAl20 (Cr-base plate with 20 Al plugs) with a purity of wCr  99.9 wt%, wAl  99.5 wt% and one MoS2 target with a purity of wMoS2  99.5 wt% were used. The MoS2 target was mounted on one of the four dcMS cathodes. In order to ensure a sufficiently high adhesion strength between the top layer and the substrate, a (Cr,Al) bond coat and a (Cr,Al)N interlayer were deposited first. The process parameters for the deposited coatings are listed in Table 1. The process parameters of the (Cr,Al)N + Mo:S top layer were explained in more detail in Bobzin et al. [21].

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Table 1 Process parameters for deposition of the (Cr,Al)N + Mo:S-coatings Process parameter

Unit

Deposition time, t

Min

107

Argon flux, j(Ar)

sccm

200

Nitrogen flux, j(N2)

sccm

Pressure controlled

Pressure, p

mPa

710

dcMs cathode power, PdcMS _MoS 7

kW

2

dcMs cathode power, PdcMs-crAi20

kW

3

HPPMS mean cathode power, PHPPMS-CrAl20

kW

5

Pulse length, ton

us

40

Frequency fpulse

Hz

500

Bias voltage, UBias

V

100

Analysis of the Influence of Shot Peening Parameters on Evolving Surface Textures Shot peening by steel casks is commonly used as workpiece preparation in industrial cold forging processes [20]. The surface modification supports a friction reduction in combination with the use of a lubricant. In dry metal forming, as a result of the missing lubricant, a higher friction is expected. In order to get the most suitable surface texture for a friction reduction in cold forging, different peening media were investigated [20]. Besides steel casks, ceramic beads and corundum particles were used with different peening parameters, see Fig. 2 and Table 2. The best peening parameters for a reduction of the frictional shear stress in tribometer tests are the ones shown in Table 3 [22]. Furthermore, a rather sharp and hard surface texture as the one shot by corundum particles, reduced the frictional shear stress the most. Table 2 Shot peening parameters in dependency of shot media, (Steel  St, Ceramic  Ce, Corundum  Co) Medium

Particle form

Particle size (µm)

Density (g/cm3 )

Hardness (HV)

St

Round

700–1,000

7.8

390–535

Ce

Round

125–250

3.8

~1,200

Co

Sharp

425–600

3.9–4.1

~2,600

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Ref

St

100 µm

100 µm Co

Ce

100 µm

100 µm

Fig. 2 SEM images of reference surface (Ref) and textured surfaces (St  Steel casks, Ce  Ceramic beads, Co  Corundum particles) Table 3 Resulting surface integrity after shot peening (Steel  St, Ceramic  Ce, Corundum  Co, Ref  Reference)

Surface

Peening parameters

Surface integrity

P (bar)

O (%)

HV10 (HV)

St

3

100

Ce

2

100

Co

4

100

Ref

Sa (µm)

Sz (µm)

280

2.23

25.7

410

1.38

26.8

260

4.09

59.7

200

1.63

29.5

Development and Use of a Pin-on-Cylinder Tribometer For an explicit investigation of the influence of the surface textures on the frictional shear stress, conventional tribometers as a pin-on-disc tribometer cannot be used since textures are flattened after the first revolution of the disc. To analyze the surface textures intensively, a new frictional path per revolution has to be used. Thus, a pinon-cylinder (POC) tribometer concept was realized on a lathe, see Fig. 3. The normal load is applied by means of a hydraulic actuator [23]. Due to these investigations, results regarding the influence of the contact between self-lubricating coatings and surface textures on the friction shear stress were obtained [24]. Shot peening by corundum particles as the best treatment to reach lower frictional shear stresses in this interaction was confirmed, see Fig. 4.

Investigation of Friction and Adhesion Behavior …

1621 Pin-On-Cylinder with axial feed

Pin-On-Disc Frictional path

Pin

Pin

FN

Friction

FN ω

Feed

ω

ω Friction

Cylinder Disc

Frictional path

Frictional shear stress τ [MPa]

Fig. 3 Schematic overview of pin-on-cylinder Tribometer 1200 1000 800 600 400 200 0 Steel

Ceramic

Corundum A

Reference

B

Fig. 4 Frictional shear stress during tribometer tests with uncoated (a) and coated (b) tools

Effect of Surface Textures on the Punch Force in an Extrusion Process After investigating surface textures using tribometer tests, the surface textures were applied in an extrusion process [25]. As coating a slightly adapted one compared to the tribometer tests were used. The punch velocity was a constant vP  5 mm/s and the punch force was measured throughout the whole process. The punch force in extrusion processes is greatly influenced by the friction [26]. A lower punch force corresponds to a lower friction during the process. The results showed that the lowest punch force was achieved from the combination of ceramic shot-peened surface texture with self-lubricating tool coating, Fig. 5. The comparison between the results of the tribometer tests and the extrusion process shows a discrepancy. A hard and rough surface structure results in low frictional shear stresses in the tribometer tests, whereas a softer and slightly rough surface texture leads to lower punch forces in the extrusion processes. This lack of knowledge needs to be overcome. To do so, smoothing effects for each surface texture need to be investigated. In dependency of the contact normal force, smoothing of the different surface textures manufactured by shot peening with different media is researched.

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R. Hild et al. 1600

Punch Force F P [kN]

1400 1200 1000 800 600 400 200 0 Steel

Ceramic

Corundum A

Reference

B

Fig. 5 Maximum punch forces of workpieces with different surface structures and uncoated (a) and coated (b) tools

Experimental and Numerical Set Up This section describes the characterization evaluation methods of the surface textures using roughness and hardness measurements, the used tools to conduct the ring compression tests (RCT) as well as the numerical simulation to receive the calibration curves for the RCT. As specimen material the quenched and tempered steel 16MnCr5 (DIN: 1.7131, AISI: 5115) in an untreated state was used. These steels are used for highly loaded parts in automotive or aircraft industries, e.g. gears, spindles, rods, and rams. Therefore, the material is very important for the forming industry. As a tool material, the cold working tool steel X155CrMoV12 (DIN: 1.2379, AISI D2) was used. It is a typical tool material in cold forging [20]. X155CrMoV12 was hardened to 62 HRC and the tool surfaces were polished (Rz  1) according to the industrial standard. The experiments were conducted using an upsetting tool, see Fig. 6. In order to analyze the impact of different surface textures during the upsetting process, the force was measured during the process, as well as the height, inner diameter and outer diameter of the specimen according to Burghoff [14]. Before and afterwards the experiments, the surface roughness and hardness of the specimens were analyzed. For the experiments, the hydraulic press HPX 400 of Schuler AG was used. The hydraulic press has a maximum force of F max  4,000 kN. During the experiments a continuous velocity of vP  5 mm/s was applied.

Investigation of Friction and Adhesion Behavior …

1623

Fig. 6 Upsetting tools with positioned ring for ring compression test

Surface Properties The surface hardness was quantified by means of Vickers hardness according to DIN EN ISO 6507 and ASTM E384 using a Wilson Universal Hardness Tester UH250. 16 measurements using a testing load of 30 N were performed on polished cross sections of each specimen. The 16 measurements were then averaged. The surface roughness was measured using a combined roughness and contour measurement system Hommel-Etamic nanoscan 855 made by Jenoptik AG, Germany. According to preliminary work, the surface parameters Sa and Sz were evaluated based on standards ISO 25178 and EUR 15178N in a measurement line of 4.8 mm. A contact tip with a tip angle of 60° and a tip radius of 2 µm was used. Ra describes the averaged roughness of a surface, Sz its maximum height. Furthermore, laser scanning microscopy analysis was conducted in order to measure the size of each indentation. Laser scanning microscope Keyence VK-X100 was used. Each indentation has a specific size. In order to get the maximum resolution, for each indentation, a specific measuring area was selected, see Table 4.

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Table 4 Initial surface roughness and surface roughness after RCT Surface

Roughness Rai (µm)

Rzi (µm)

Rae (µm)

Rze (µm) 1.775

B1

1.01

6.28

0.187

B2

1.75

8.38

0.244

1.943

B3

4.11

21.50

0.221

2.620

Ref

1.55

7.54

0.258

2.426

Bb1

1.01

6.28

0.305

2.409

Bb2

1.75

8.38

0.260

2.442

Bb3

4.11

21.50

0.276

3.287

Refb

1.55

7.54

0.380

3.064

Numerical Set up and Calibration Curves To determine the calibration curves for the ring upsetting tests, a finite element simulation with Forge NxT 2.1 simulation software was set up. Two upsetting plates with a diameter of d 0  80 mm served as the basis for this. These upsetting plates were defined as rigid. The workpiece was defined with the dimensions from the experiment and defined as elasto-plastic. The cross-linking of the component was defined with an edge length of 0.25 mm, which led to 443,831 elements. The simulation was carried out step by step according to DIN. The ring was compressed by 0.4 mm from simulation to simulation, to an end height of hF  3.2 mm, see Fig. 7. The infeed was carried out by adjusting a hydraulic press with a constant speed of vP  5 mm/s. The calibrated friction curves are as a function of the friction factors of m  {0.0; 0.02; 0.03; 0.04; 0.05; 0.06; 0.08; 0.10; 0.12; 0.15; 0.2; 0.3; 0.4; 0.5; 0.577}, see Fig. 8. These are applied over the height on the abscissa and the inner diameter on the ordinate. The calibrated friction curves are extended by the experimental analysis and the respective support points are entered. Thus, the friction values for the respective surface structures and the dependency of the coating are determined.

Fig. 7 Set up for numerical ring compression test in order to distinguish calibration curves

Von Mises stress σM [MPa] 400

200

0

Investigation of Friction and Adhesion Behavior …

1625

Ring compression test Height h [mm] 6.5

m=0.12

6

5.5

5

4.5

4

m=0.577 m=0.5 m=0.4 m=0.3 m=0.2 m=0.15

3.5 4 5 6 7 8 9 10

m=0.06 m=0.05 m=0.04 m=0.03 m=0.02 m=0.0

11 12

Inner diameter d [mm]

7

13 14

Fig. 8 Calibrated friction curves for ring compression tests

Results and Discussion After the ring compression tests had been carried out, the samples were entered into the calibration curves depending on their inner diameter and height. Due to the positioning accuracy of the press, slight variations in the support points have resulted depending on the sample height. All tests performed are in a range of curves for m  0.3 to m  0.15, see Fig. 9. This clearly shows that high coefficients of friction are achieved under dry tribological conditions. Nearly all ring upsetting specimens are initially at a lower level and then rise rapidly, so that the height inside diameter values ends closer to higher coefficients of friction. The shot-peened upsetting specimen without the use of a coated upsetting plate shows this tendency most clearly. Initially, the sample has values that correlate with those of m  0.1. The next support point refers to the curve with the coefficient of friction m  0.2, until finally, a further increase occurs and the specimen achieves the highest coefficient of friction of m  0.3. In contrast, the steel-blasted specimen in combination with a coated upsetting plate shows a relatively constant course and the lowest coefficient of friction. This is located almost constantly on the curve of m  0.15. The ceramic-blasted upset sample, which was upset without a coated upsetting plate, has a friction value of about m  0.18. The equally treated upsetting sample using a coated upsetting plate achieves a friction value of m  0.16. This surface structure thus generally achieves low friction values, which, however, do not become as low as the combination of a steel-blasted sample with a coated upsetting plate. Samples blasted with aluminium oxide generally have a higher coefficient of friction. With an uncoated upsetting plate, this is at m  0.2 and with a coated upsetting plate at about m  0.22. At the beginning of the upsetting process, the friction of the upsetting sample blasted with white corundum using an uncoated upsetting plate was about

1626

R. Hild et al. Height h [mm] 6.5

6

5.5

5

4.5

4

3.5 4 5 6 7 8 9 10 11 12 13 14

BW1

BW2

BW3

BWb1

BWb2

Inner diameter d [mm]

7

BWb3

Fig. 9 Determination of friction values by ring compression tests and calibration curves

m  0.26. In general, it can therefore be seen that the self-lubricating coating of (Cr,Al)N + WSx reduces friction. This can vary in intensity in some cases. At the same time, the influence of the surface structure of the semi-finished product on friction is demonstrated. The partially significant change in the friction values over the course of the compression also shows that there is a continuous change in the tribological system. At the beginning of the compression, the surface structure is completely preserved. Thus, lower coefficients of friction are achieved at the beginning than in the further course of the tests. The reduction of the contact surface thus results in a reduction of the coefficient of friction up to complete levelling. This correlates with observations during compression. During the intermittent removal of the upsetting samples, it was determined that the surface structure could still be partially visually detected up to a height of about 5.5 mm. After further compression, the initial structure was completely flattened. Furthermore, it is shown that the roughest and hardest surface structure, which is produced by white corundum, achieves the highest coefficients of friction. Samples blasted with ceramic achieve the continuously lowest coefficient of friction. Only the upsetting specimen blasted with steel using a coated upsetting plate achieved slightly lower friction values. This proves that for an optimized friction condition, a slight roughening of the surface with simultaneous increased hardening of the edge zone is expedient. The cross section of the ring upsetting specimens reflects this. According to Sofuoglu et al. [16], a double convex upsetting sample has a higher friction factor than a convex–concave upsetting sample, see Fig. 10. This image is most clearly visible in sample BW3. Both the inner diameter and the outer diameter are clearly bent outwards. Sample BWb2, on the other hand, has an almost straight inner diameter, which also proves that these samples have the lowest coefficient of friction.

Investigation of Friction and Adhesion Behavior … Ceramic shot peened

Steel shot peened

Corundum shot peened

BW1

BW2

BW3

BWb2

BWb3

hf

1627 Reference surface texture

df

BWb1

Fig. 10 Visual analyzation of ring compression specimens after testing and comparison to literature

Summary and Outlook The resulting discrepancy between the frictional shear stress during tribometer testing and the punch force in an extrusion process led to an investigation of the plastic deformation of different surface textures. A reference surface texture machined by turning as well as several shot-peened specimens, peened by steel, and by ceramic casks and by corundum particles was investigated. Due to different pretreatments, different surface integrities evolved. These surface textures were tested with increasing normal contact forces. Rising contact normal forces resulted in larger indentation areas. Furthermore, a harder and rather flat surface texture resulted in the smallest indentation area, whereas a flat but soft surface texture resulted in the greatest indentation area. Respective to the lowest punch force, an ideal workpiece has a hard surface texture with a low roughness. In order to separate the effects of roughness and hardness of the surface texture, further researches have to be done. Hardened and unhardened with an identic roughness as well as rougher or flatter surface textures with an identic hardness need to be investigated. Acknowledgements The research was funded by the German Research Foundation (Deutsche Forschungsgemeinschaft DFG) within the priority program “Dry metal forming—sustainable production through dry processing in metal forming” (SPP 1676). The authors also thank Presswerk Krefeld GmbH for their support supplying tools and workpieces as well as conducting the full forward extrusion process.

References 1. Bay N, Azushima A, Groche P, Ishibashi I, Merklein M, Morishita M, Nakamura T, Schmid S, Yoshida M (2010) Environmentally benign tribo-systems for metal forming. Ann CIRP 59(2):760–780 2. Vollertsen F, Schmidt F (2014) Dry metal forming: definition, chances and challenges. Int J Precis Eng Manuf Green Technol 1/1:59–62 3. Vollertsen F, Schmidt F (2014) Dry metal forming: definition, chances and challenges. Int J Precis Eng Manuf Green Technol 1(1):59–62 4. Müller K (2003) Praktische Oberflächentechnik – Vorbehandeln – beschichten – Beschichtungsfehler – Umweltschutz; JOT Fachbuch

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5. Teller M, Bambach M, Hirt G (2015) A compression-torsion-wear-test achieving contact pressures of up to eight times the initial flow stress of soft aluminium. CIRP Ann Manuf Technol 64:280–292 6. Teller M, Bambach M, Hirt G, Ross I, Temmler A, Poprawe R, Bolvardi H, Prünte S, Schneider JM (2015) Investigation of the suitability of surface treatments for dry cold extrusion processoriented tribological testing. Key Eng Mater 651–653:473–479 7. Murakawa M, Koga N, Kumagai T (1995) Deep-drawing of aluminum sheets without lubricant by use of diamond-like carbon coated dies. Surf Coat Techol 76:553–558 8. Murakawa M, Takeuchi S (2003) Evaluation of tribological properties of DLC films used in sheet forming of aluminum sheet. Surf Coat Techol 163:561–565 9. Osakada K, Matsumoto R (2000) Fundamental study of dry metal forming with coated tools. CIRP Ann Manuf Technol 49(1):161–164 10. Kataoka S, Motoi A (2005) Improvement in DLC thin film adhesion and its application to dry deep drawing. J Jpn Soc Technol Plast 46(532):412–416 11. Tamaoki K, Kataoka S (2008) Study of deep drawing using diamond coated tools. J Mater Test Res Assoc Jpn 53(4):247–253 12. Tamaoki K, Kataoka S, Minamoto K (2007) Dry deep-drawing with use of electroconductive ceramic tools. Proc Int Conf Trib Manuf Pro (2007) 175–179 13. DIN50323-2 (1995) Tribologie, Verschleiß 14. Burghoff M (1967) Über die Ermittlung des Reibwertes für Verfahren der Massivumformung durch den Ringstauchversuch. Industrie-Anzeiger Werkzeugmaschine und Fertigungstechnik 15. Doege E, Behrens B-A (2010) Handbuch Umformtechnik. Springer, Berlin Heidelberg, Berlin, Heidelberg 16. Sofuoglu H, Hasan G, Rasty J (2001) Determination of friction coefficient by employing the ring compression test. Trans ASME 17. Deghan M, Fathallah Q, Gerdooei M, Doai J (2013) Analysis of ring compression test for determination of friction circumstances in forging process. Appl Mech Mater 249–250:663–666 18. Bay N, Petersen SB, Martins PAF (1998) An alternative ring-test geometry for the evaluation of friction under low normal pressure. J Mater Process Technol 79:14–24 19. Mattfeld P (2014) Tribologie der zinkphosphatfreien Kaltmassivumformung. Dissertation 20. Groche P, Stahlmann J, Hartel J, Köhler M (2009) Hydrodynamic effects of macroscopic deterministic surface structures in cold forging processes. Tribol Int 42 21. Klocke F, Trauth D, Schongen F, Shirobokov A (2014) Analysis of friction between stainless steel sheets and machined hammer peened structured tool surfaces. Prod Eng Res Dev 8:263–272 22. Bobzin K, Brögelmann T, Bastürk S, Klocke F, Mattfeld P, Trauth D (2015) Development of an insitu plasma treatment of X155CrMoV12 for a (Cr, Al)N PVD tool coating for dry metal forming in cold forging. Dry Metal Form Open Access J 1(1):57–62 23. Bobzin K, Brögelmann NC, Bastürk S, Klocke F, Mattfeld P, Trauth D (2015) Tribological behavior of (Cr1-x Alx )N/WSy PVD tool coatings for the application in dry cold forging of steel. Dry Metal Form Open Access J 1(1):152–158 24. Trauth D, Hild R, Mattfeld P, Bastürk S, Brögelmann T, Bobzin K, Klocke F (2016) Advances in dry metal forming of low alloyed steels for cold forging using a (Cr,Al)N tool coating and surface structures on workpieces. In: Proceedings of the 12th international conference a coatings in manufacturing engineering, Hannover, Germany, 31 March–1 April 2016 25. Bobzin K, Kruppe NC, Arghavani M, Hoffmann DC, Klocke F, Mattfeld P, Trauth D, Hild R (2017) Mechanical and tribological characterization of self-lubricating (Cr1-x Alx )N coatings for deposition on complex-shaped forging tools. Mechanical and tribological characterization of self-lubricating (Cr1-x Alx )N coatings for deposition on complex-shaped forging tools. Dry Metal Form Open Access J 3:81–89 26. Bobzin K, Kruppe NC, Arghavani M, Hoffmann DC, Klocke F, Mattfeld P, Trauth D, Hild R (2018) Einfluss von Oberflächenstrukturierungen auf die Stempelkraft beim Vollvorwärtsfließpressen von 16MnCr5. Dry Metal Form Open Access J (4):25–30

Effects of Emissivity on Combustion Behavior of Energetic Materials Elbert Caravaca, David Bird, Henry Grau, Viral Panchal and Nuggehalli M. Ravindra

Abstract Emissivity is defined as the ratio of the energy radiated from a material’s surface to that radiated from a blackbody (a perfect emitter) at the same temperature and wavelength and under the same viewing conditions. It is a dimensionless number between 0 (for a perfect reflector) and 1 (for a perfect emitter). Knowledge of surface emissivity is important both for accurate non-contact temperature measurement and for heat transfer calculations [1]. Piobert’s Law states that all burning occurs at the surface layer by layer and the burning is normal to the surface [2]. This may cause changes in the burn surface response depending on the surface. Another important characteristic of combustion is Saint Robert’s Law (a.k.a. Vieille’s Law) which explains the effect of pressure [3]. This paper will investigate through literature and propose a technical approach. The goal will be to assess the effect of a materials emissivity to its combustion behavior and identify a correlation if one exists. As stated above, emissivity does influence heat transfer which is also a contributor to combustion behavior. Keywords Emissivity · Heat transfer · Combustion behavior · Energetic materials · Nitrocellulose · Celluloid

Introduction The Department of Defense (DoD), and more specifically the US Army, has interest in developing novel methods of manipulating combustion behavior in legacy materials which utilize nitrocellulose as the main energetic binder. Nitrocellulose is an inexpensive and mass-produced commodity highly utilized in gun propellant manufacture for the DoD and commercially. Celluloid materials are being investigated E. Caravaca (B) · D. Bird · H. Grau · V. Panchal Propulsion Technology and Prototyping Division, US Army ARDEC, Picatinny, NJ 07806, USA e-mail: [email protected] N. M. Ravindra New Jersey Institute of Technology, Newark, NJ 07102, USA © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_152

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at Picatinny Arsenal as a combustible cartridge case material. The combustible cartridge cases are used in large caliber weapon systems for easy transport and utility of gun propulsion charges during gun firing. The sole purpose of these cases is to protect the energetic charges and be completely consumed during the interior ballistic cycle in any weapon platform. Celluloid is composed of nitrocellulose and camphor and has very low burn rates. Celluloid also generates soot which is very undesirable. Soot has an emissivity value of 0.7 [4]. The emissivity of celluloid is currently unknown. Foamed celluloid, on the other hand, has a much higher burn rate, 89 cm/s [5]. Current production combustible cases manufactured by Esterline Defense Group in Coachella, CA, have similar burn rates but are manufactured using a wet mold process [6]. Celluloid is a polymer-like material and has the potential to incorporate material additives to influence combustion rates. Minimizing inert or low energy additives reduces the risk of unwanted combustion residue. There is a counterintuitive engineering challenge with combustible materials. Combustible cases (typically weak) are subject to the fact that increasing strength must not be at the expense of maintaining clean combustion. But improving combustion behavior through material solutions would enable using additives for increased strength. The goal is to understand the effect of modifying the radiant properties of celluloid and the extent to which it can be used to improve combustion behavior. The effort will start with baseline characterization of celluloid and corresponding emissivity value. Then, select the appropriate additives or treatments to modify the emissivity values of the celluloid samples. After samples are prepared, the new emissivity values will be obtained. Finally, the samples will be tested for burn rate.

Literature Review Emissivity Relative to Energetic Materials During the literature search, some papers were found to show some connection with shock-induced chemical reactions relative to the material emissivity. Basset and Dlotta found the data suggested that there are some means of effecting or disrupting the shock to deflagration and shock to detonation transitions of powdered HMX. Spectral data was collected during flyer plate experiments in order to track the change in hot spots and how that tracks with HMX detonation/deflagration output [7]. The relevance of this data analysis implies that emissivity can be a means to understand what is happening with energetic materials at or near the surface. Figure 1 shows the various shock velocities through the HMX samples and the interesting trends. The temperature rise was always followed by an increase in emissivity and this occurred twice for each sample. Although very interesting results obtained by the authors also identified that the two temperature spikes most likely are due to compactification of the low-density powders. Typical explosives are highly densified to maximize effectiveness as well

Effects of Emissivity on Combustion Behavior of Energetic …

1631

Fig. 1 Results of flyer plate experiments and the corresponding emissivity results [7]

as avoid safety issues with adiabatic compression of voids in the materials [8]. With voids comes the concern that, under rapid shock loading, the air is rapidly compressed and the temperature rises which can kick start the chemical reaction kinetics to deflagrate/detonate the HMX.

Influencing Emissivity of Polymeric Materials Another aspect to exploring the concept of effecting combustion behavior of combustible materials is the method of changing emissivity. Cernuschi et al. investigated the effect of doping various polymers and the change in emissivity. It is clear that various doping levels of materials such as polyaniline, polythophene and polypyrrolle have an impact on emissivity. All polymers were doped with NH3 and all exhibited a decrease in emissivity from partially to completely doped samples. Below are tables with the results of these measurements [9] (Tables 1, 2 and 3). The above work leads to the derivation of the equation, Fig. 2, as a means to address doping level and the change in emissivity. The authors identified the sensitivity of the measured emissivity and “true” emissivity. The same paper provided a graphical depiction of the discrepancy between measured and actual emissivity using the value of a true black body reference. The reason for this is that better results are generally obtained using the direct technique with a high-temperature difference between sample “S” and background “B” and a very

1632 Table 1 Data of emissivity values of polythiophene [9]

Table 2 Data of emissivity values of polypyrrole [9]

Table 3 Data of emissivity values of polyaniline [9]

E. Caravaca et al.

Sample

Sample temperature (°C)

Emissivity

Climatic chamber average temperature (°C)

Completely undoped polytiophene

20

0.932

16.5

25

0.939

30

0.944

Partially doped polytiophene

20

0.906

25

0.917

30

0.921

Doped polytiophene

20

0.760

25

0.770

30

0.761

Sample

Sample temperature (°C)

Emissivity

Climatic chamber average temperature (°C)

Partially doped polypirrole

20

0.847

13.5

25

0.851

30

0.857

Doped polypirrole

20

0.817

25

0.825

30

0.824

Sample

Sample temperature (°C)

Emissivity

Climatic chamber average temperature (°C)

Completely undoped polyaniline

20

0.913

9.5

25

0.915

30

0.916

Partially doped polyaniline

20

0.888

25

0.891

30

0.889

Doped polyaniline

20

0.725

25

0.741

30

0.733

16.5

16.5

13.5

9.5

9.5

Effects of Emissivity on Combustion Behavior of Energetic …

1633

Fig. 2 Derived relationship with emissivity and differences in obtained temperatures [9]

Table 4 Correlation with the amount of additive to an LDPE on time to ignition or TTI [9]

Formulations

Aluminum particles

Grade content (wt%)

Film thickness (μm)

1A20-65

A20

1

65

1A20-100

A20

1

100

3A20-100

A20

3

100

5A20-100

A20

5

100

5A40-100

A40

5

100

5A60-100

A60

5

100

uniform background surrounding the sample and the mirror. These factors will need to be considered for the work proposed in this study. The relevance of the above work is the fact that celluloid is a polymer like material and could be doped with some materials that can increase or decrease emissivity. This will allow for experimentation of combustion behavior relative to these changes. The technique of obtaining emissivity will need to be identified in order to provide useful data when relating to combustion rates. One major concern is the relatively high values of the polymers reported in the paper [9]. If increasing emissivity increases combustion rates, a very narrow window could be used to influence burn rates. This will have to be determined experimentally in order to identify the sensitivity of emissivity to burn rate. Another study was found to show a correlation with the amount of additive to an LDPE on time to ignition or TTI. Initially, the LDPE was coated with varying amounts of Al coating. The sample numbers and their corresponding details are presented in Figs. 3, 4 and Table 4. Cernusch et al. found a connection with changes in emissivity and TTI of the LDPE samples. This does indicate the promise of influencing the burn rates of combustible materials. Figure 4 shows the effect more clearly. The use of varied particle-sized Al has a significant effect on emissivity. When comparing the plot and the emissivity values, it is clear that there is a correlation as summarized in Table 5. One concern is that when heat flux is increased, the effect is less prominent. This may mask the effect of emissivity with celluloid. Figure 5 displays this change with increased heat flux. The authors also postulated that if direct flame contact occurs, this effect may not be realized. They reference the mirror effect with regards to radiant heating; it will

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Fig. 3 Correlation with the amount of additive to an LDPE on time to ignition or TTI [9]

Fig. 4 Effect of particles to TTI [9]

Effects of Emissivity on Combustion Behavior of Energetic …

1635

Table 5 Emissivity values of various coatings [9] Coatings

ρ1

t1

ε1

1A20-65

0.231 ± 0.005

0.260 ± 0.012

0.509 ± 0.013

1A20-100

0.245 ± 0.006

0.192 ± 0.012

0.563 ± 0.013

3A20-100

0.349 ± 0.001

0.110 ± 0.003

0.541 ± 0.004

5A20-100

0.369 ± 0.006

0

0.631 ± 0.006

5A40-100

0.372 ± 0.009

0.132 ± 0.040

0.497 ± 0.041

5A60-100

0.370 ± 0.002

0.156 ± 0.006

0.475 ± 0.007

Fig. 5 Effect of heat flux related to increased A20% content and TTI [9]

lose its impact if direct flame is applied. In large caliber, combustion events within an interior ballistic cycle of a weapon flame and hot gases are present and depend on the events and where this emissivity correlation may be masked. This will be addressed experimentally. As described in this paper, there are various methods to influence emissivity. After the celluloid materials are characterized for the baseline emissivity values, further investigation will be conducted. This will help select the best candidates to incorporate these additives or modifications to the celluloid materials.

Methods of Measuring Emissivity Celluloid is processed in various ways. The methods that will be utilized for this study will use the method described in Fig. 6. Celluloid is processed in a horizontal mixer

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Fig. 6 The celluloid block process method [4]

with solvents. It is composed of ~11.0% nitrated nitrocellulose at ~80% by weight and camphor at 20% by weight. Commercial manufacture of celluloid typically involves mixing nitrocellulose and camphor in the presence of solvents, such as ethanol and acetone. A common celluloid manufacturing process, known as “blocking” involves mixing the nitrocellulose, camphor, and other ingredients, followed by straining, roll milling and “hiding”. A selected number of “hides” are then blocked at a desired pressure and temperature into a fused block, which is then sliced into sheets of desirable thickness after a conditioning period. Alternatively, celluloid can be manufactured by “film casting.” which involves mixing nitrocellulose, camphor, and other ingredients, and subsequently casting, and drying the mixture into film sheets of a desired thickness. Celluloid, when molded or cast, is translucent and is amber in color. Based on the definition of emissivity, celluloid may be defined as a grey body material [10]. This data will be generated using ATR-IR spectroscopy at ARDEC, Picatinny, NJ to support this effort. A subsequent study will attempt to correlate emissivity and combustion. Okada et al., utilized ATR-IR spectroscopy to determine the spectral response and compared it to black body curves of low-density polyethylene (LDPE, Showa Denko Co.), high-density polyethylene (HDPE, Showa Denko Co.), nylon 6 (Toray Co.) pellets, polyethylene terephthalate films (PET, Toray Co.), and polyethylene foam (PE, Asahi Kasei Co.), and nylon 6 fabric (Toray Co.) materials [11]. Spectral measurements will be performed and then the calculations identified in the Okada paper will be used to determine emissivity values. Table 6 reported by Okada shows good agreement with calculated emissivity and measured emissivity using ATR-IR method. One concern in using the ATR-IR method is the need for using non-translucent materials as a potential criteria for this method. Again, through experimentation, we will discover the range of emissivity of the proposed material celluloid [12].

Effects of Emissivity on Combustion Behavior of Energetic … Table 6 Obtained emissivity values of various materials [9]

1637

Sample

Emissivity εATR (/100 μm)

Emissivity ε [6] (/100 μm)

Penetration depth d p (μm)

LDPE

0.27

≈0.30

0.5–3.1

HDPE

0.28

≈0.30

0.5–3.1

Nylon 6

0.73

≈0.75

0.5–3.6

PET

0.81

≈0.80

0.5–3.6

PI

0.83

–

0.8–5.5

PE foam

0.23

–

0.5–3.1

Nylon fiber

0.61

–

0.5–3.6

Methods of Measuring Burn Rate After the samples are prepared, each sample will be subjected to low pressure combustion testing. A Design Integrated Technology (DIT) Strand Burner System will be used to characterize the samples. The system at ARDEC Picatinny can be

Fig. 7 Picture of control console [13]

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Fig. 8 Test cell housing and temperature controller [13]

tested from 500 to ~3000 psi. Temperature can be controlled during the test. The samples will be ~5.5 in. in length and ~0.5 in. in width. The thickness can be ~0.5 in. A standard inhibitor will be applied to the samples except for the burn surface in question. Figures 7 through 10 show the test equipment that will be used. The samples are fixed to break wires which are wrapped around the pair of nuts which complete a circuit. Three sets of nuts are present. Two are for calculating the burn rate, and the third to ignite the sample through resistive heating. The included data acquisition system captures the burn velocity and corresponding changes in pressure and temperature. A burn rate will be provided as the results.

Summary and Path Forward The effect of emissivity on combustion behavior of relevant energetic materials are being investigated. There is some indication of a direct or indirect impact on combustion behavior related to the emissivity of materials. Some work has been done to correlate the emissivity of materials to chemical reaction kinetics. Time to ignition is also shown to be influenced by emissivity of the material. Celluloid is the material

Effects of Emissivity on Combustion Behavior of Energetic …

Fig. 9 Test chamber for strand burner sample holder [13]

Fig. 10 Strand Sample holder [13]

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that is becoming widely used at ARDEC for combustible case application in large caliber weapon systems. Influencing combustion behavior through non-traditional means will enable novel material solutions and design ideas for celluloid in other weapon platforms. This effort will investigate characterizing celluloid materials before and after modification. A set of emissivity modifying materials will be selected and subsequent mixes utilized to prepare various samples. The emissivity values will be obtained for these sets of samples. ATR-IR will be used at ARDEC to analyze the samples. The materials will then be prepared for combustion testing. The burn rates and emissivity values will be compared and any trends will be identified. Acknowledgements ARDEC—Dana Kaminsky, Duncan Park, Joseph Laquidara, Carlton Adam, David Keyser, Scott McDonald, Keith Luhmann. ARL—Richard Beyer John Schmidt, Paul Conroy. Family—Rezeile, Elyssa, Alvaro and Lianna Caravaca for supporting the primary author getting this paper and work started.

References 1. (2014) What is emissivity and why is it important? In: NPL. http://www.npl.co.uk/reference/ faqs/what-is-emissivity-and-why-is-it-important-(faq-thermal). Accessed 30 April 2018 2. (2018) Piobert’s law. In: Wikipedia. https://en.wikipedia.org/wiki/Piobert%27s_law. Accessed 30 April 2018 3. Nakka R (2003) Solid propellant burn rate. In: Richard Nakka’s experimental web site. https:// www.nakka-rocketry.net/burnrate.html. Accessed 30 April 2018 4. Young M (2013) Foamed celluloid combustible material. US Patent 8597444B1, 3 Dec 2013 5. (2018) Emissivity table. In: Thermoworks. https://www.thermoworks.com/emissivity_table. Accessed 30 April 2018 6. (2017) Armtec products for military applications. In: Esterline defense technologies. http://www.esterline.com/defensetechnologies/CombustibleOrdnance/TankCases.aspx. Accessed 30 April 2018 7. Basset WP, Dlotta DD (2016) Shock initiation of explosives: temperature spikes and growth spurts. Appl Phys Lett 109:091903 8. Wolfs F (2013) Heat and the first law of thermodynamics. http://teacher.pas.rochester.edu/ phy121/lecturenotes/Chapter17/Chapter17.html. Accessed 30 April 2018 9. Cernuschi F, Russo A, Piana GM, Mutti P, Viviani L (1997) Emissivity measurements at room temperature on polymeric and inorganic samples, QIRT. https://doi.org/10.21611/qirt.1996. 007 10. Infrared basics infrared energy, emissivity, reflection & transmission, Williamson Corporation, 70 Domino Drive, Concord, MA 01742; E-Mail: [email protected] 11. Okada T, Ryohei I, Ando S (2016) Analysis of thermal radiation properties of polyimide and polymeric materials based on ATR-IR spectroscopy. Department of Chemical Science and Engineering. J Photopolym Sci Technol 29(2):251–254

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12. Sonnier R, Laurent F, Gallard B, Boudenne A, Lavaud F (2015) Controlled emissivity coatings to delay ignition of polyethylene. Materials 8:6935–6949. https://doi.org/10.3390/ma8105349 13. (2016) Strand burner system. In: Design integrated technology. http://ditusa.com/strand_burner.php?gclid=Cj0KCQjw_ZrXBRDXARIsAA8KauSdcIJb3f deuPtUZ5OjswKFOP9XQ5Y7oM5DMNkaAFGYey8JHlgUumhQaAm2MEALw_wcB. Accessed 30 April 2018

Self-healing in Materials: An Overview Samiha Hossain and Nuggehalli M. Ravindra

Abstract Conventional materials often fail in various ways even from relatively lowimpact cyclic forces. In order to combat this phenomenon, smart materials have been developed. They exhibit adaptive capabilities to external stimuli, such as physical, mechanical or chemical changes in their environment. This makes them a topic of immense interest as they have many applications of importance in the scientific field. In order to make advancements of significance in this arena, a better understanding of the fundamentals is required. This overview attempts to summarize the self-healing mechanisms in a wide variety of materials including natural composites, polymers, metals and ceramics. Keywords Self-healing · Polymeric materials · Autonomous self-healing · Induced self-healing

Introduction Self-healing Concept A wide variety of materials are available for multiple applications in the engineering field; however, a very few are practical. All materials are exposed to multiple external stimuli such as thermal, mechanical, chemical stresses or radiation, which make them susceptible to damage. This can lead to structural degradation and internal crack formation, which alter the physical and mechanical properties of the material. When structural materials are impaired, there are very few ways the “device” can be conserved so that its functional lifetime can be extended. Ideally, the material would repair itself by continuously sensing external stimuli and respond to the damage S. Hossain (B) · N. M. Ravindra Interdisciplinary Program in Materials Science & Engineering, New Jersey Institute of Technology, Newark, NJ, USA e-mail: [email protected] © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_153

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Fig. 1 Lifetime extension of engineered materials by the implementation of self-healing principle [3]

caused accordingly. This makes the material durable, more reliable, easier to maintain and reduce repair costs. Therefore, there is a need to develop “smart” materials [1]. Conventional methods of repairing materials include welding (reformation of surface bonds at site of fracture by heat, infrared or other external stimulation) or patching (replacement or covering/coating with different material at site of fracture). However, these methods have inadequate reliability as the location of repair is still the weakest point in the material thus making it likely to fail again. Materials with the ability to heal autonomously and consistently at the molecular level without the need for external intervention are thus required to mitigate these issues [1]. The ideal self-healing material should have several characteristics that make it technologically and economically appealing. S. van der Zwaag defined the terms for the ideal material and a minimal self-healing material. The ideal material should be able to heal damage of any size completely and limitlessly without the need for external intervention (autonomously). The healed site should have equal or superior properties to the original matrix and should be cost-effective. A minimal self-healing material (a more realistic situation) is able to heal small instances of damage partially and at least once. It might need external stimulation to induce the healing process and the healed site has properties that are inferior to host matrix and the material is not cost-effective [2]. In practice, the performance of healed materials decreases with time as demonstrated in Fig. 1. In the last several years, various methods have been developed to produce selfhealing materials but can overall be grouped into two categories: induced and autonomous healing. This study focuses on the physical and chemical mechanisms employed in self-healing materials.

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Induced (Non-autonomous) Self-healing Materials with self-healing properties that can be triggered by an external stimulus such as light or heat undergo induced self-healing which involves molecular interdiffusion occurring at the fracture interface which facilitates the reformation of bonds [1].

Autonomous Self-healing There is two types of autonomous self-healing: capsule based and channel based. The overall idea is to take advantage of structural modifications such as microcapsules or fibers to induce a structural healing in the material. These two approaches use the same process of embedding a healing agent (such as an appropriate catalyst or initiator) into a membrane that separates the agent from the polymeric matrix. Capsule-based approaches make use of the dispersion of the catalyst in the polymeric matrix and healing agent encapsulated into dispersed microspheres. Channelbased approaches use long fibers organized in a 3D arrangement throughout the matrix. Instead of taking advantage of hollow spheres, in this case, the healing agent is enclosed in hollow channels or fibers that are organized in a network system. This method has the advantage over the capsule-based one that once the crack has propagated through the structure, with consequent healant depletion in the area, new healant can be supplied to the fracture site through continuous distribution via micro-channels. These two approaches commonly do not require external stimuli to be activated. Figures 2 and 3 illustrate the self-healing process.

Fig. 2 Induced self-healing requires an external stimulus to trigger the process (adapted from Blaiszik et al. [4])

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Fig. 3 Autonomous self-healing does not require external stimuli as the healant is contained within capsules or vascular systems (adapted from Blaiszik [4])

Natural Composites Most biological materials exhibit unique properties, which make them the inspiration for many of the current synthetic materials. One of the most impressive features is their ability to organize their structure according to their scale, which defines their mechanical properties among many. Currently, no synthetic material comes close to the complexity and efficiency of natural composites. Bone and nacre are two of the most studied structures and serve as a standard to emulate.

Bone Bone is the main structural component in most vertebrates. It is mostly composed of collagen fibers and inorganic bone mineral called Dahllite in the form of small crystals but depending on its function and structure, it can have differing organizational types and structures and can be made of a variety of proteins. For the purpose of this study, the main focus is on the hierarchical organization of the type of bone composed of a cortical bone structure and a cancellous bone (or spongy) structure. These structures optimize the bone’s function which includes self-regeneration, structural and mechanical properties, etc. [1]. Figure 4 shows the structure of the bone and its hierarchical architecture from the nanoscale to the macroscale. At the nanoscale, a brick and mortar structure is observed that is composed of thin “gluey” interfacial layers of non-collagenous proteins holding together the mineralized collagen fibrils [1]. Upon fracture, the gluey layers break and reform into “bridges and hidden lengths” and are able to keep the structure together despite the separation of the collagen fibers. This process highlights an intrinsic self-healing mechanism that is independent of the basic cellular regeneration. Fantner et al. hypothesized that the formation of polymerized bridges

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Fig. 4 Hierarchical levels of bone structure at different scales from macroscopic to microscopic [1] Fig. 5 Toughening mechanisms in bone (adapted from Launey et al. [7])

within the gluey layer was between either negative groups on the polymer chains or between polymer groups and the hydroxyapatite platelets [5]. Organic natural composites exhibit superior strength and toughness than their inorganic counterparts due to the anisotropic organization of the structures. They give the material properties that vary with the direction of load application. Bone is able to withstand highest compressive forces (load bearing) in the longitudinal direction [6]. It has many intrinsic and extrinsic toughening and strengthening mechanisms as shown in Fig. 5. Crack propagation is dealt with by osteons positioned perpendicular to the direction of extension of the crack which stimulates the formation of microcracks and the presence of collagen fibrils forms bridges behind the crack tip resulting in higher toughness [1, 7].

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Fig. 6 a Brick and mortar structure of abalone shell; b side view of nacre microstructure; c Voronoishaped polygonal architecture found in each nacreous layer; d staggered platelet configuration in nacre (adapted from Tran et al. [9])

Nacre Nacre or mother of pearl is one of the several different kinds of materials that compose shells of mollusks and is characterized by distinct optical properties (iridescence) due to its internal structure. It is composed of a two-phase composite of calcium carbonate (as calcite or aragonite) with an organic matrix containing glycine- and alanine-rich protein and polysaccharide [8]. At the nanoscale, a brick and mortar structure is observed that is composed of polygonal Aragonite plates glued together by thin layers of highly cross-linked organic polymer matrix (polysaccharide βchitin) [8]. The bricks are also held together by mineral bridges that provide further strength to the structure. In its natural environment, nacre is meant to keep the shell from breaking by absorbing fracture energy during impact from predator attacks. The brick and mortar structure deflects most of the crack propagation, but at the nanoscale, the bricks provide a rough surface that increases the sliding friction thus minimizing the sliding of the tiles by dissipating energy and increases the toughness of the material [8]. Figure 6 represents the structure of nacre.

Thermodynamics of Self-healing Self-healing, in one sense, is the opposite of degradation processes such as wear, fatigue, creep, etc. Most of the processes involve dissipation of energy which is irreversible and in thermodynamics, this implies the contribution to entropy [10]. Entropy, S, is a measure of irreversibility and is defined as follows: dS  (dQ/T)

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where T is absolute temperature and Q is heat. S is a result of the summation of entropies of all parts of the system. This relationship defines the second law of thermodynamics which states that the net entropy of a closed system remains constant (if reversible) or grows (if irreversible) [10]. Boltzmann defined entropy as a function of the number of microstates in a system. The equation is as follows: S  (kln ) where k is Boltzmann’s constant and  is the number of microstates corresponding to given macrostate. Microstates are arrangements of energy and matter in a system that are distinguishable at the atomic or molecular level, but are indistinguishable at the macroscopic level [10]. The more disordered a macrostate is, the higher the number of microstates within the system [10]. The mechanisms that cause degradation involve interactions on different scales of length (macro, micro and nano). Thus, in certain cases, entropy caused at a certain scale level can be compensated by entropy consumption in another level [10]. These mechanisms are independent of each other and as entropy is an additive function, the total entropy of the system can be represented as follows: Snet  Smacr o + Smacr o + Snano where each variable represents the “net”, “macro”, “micro”, and “nano” scale entropy at corresponding scales. If we consider a solid homogenous body, a perfect single crystal system would have lower microscale entropy than one with deficiencies such as grains, defects, and dislocations on the microscale. Larger scale defects such as cracks and voids contribute to entropy on the macroscale [10]. Thus, a material with a more ordered structure at the microscale has lower microscale entropy (Smicr o ) than a material with a disordered microstructure. This can be utilized to heal defects such as cracks and voids on the macroscale level by triggering healing in the microscale structures, for example, by releasing microcapsules [10]. The fracture of microcapsules increase disorder in the microscale and thus increases the entropy making it higher than  Smacr o . In this instance, the healing is done by decreasing macroscale entropy at the cost of the microscale component [10].

Polymers Polymeric materials have unique mechanical and chemical characteristics which make them suitable to work in situations where self-healing is a requirement. They are low cost, high-performance materials that has resulted in much research invested in them. Polymers can be separated into two different groups depending on the behavior they exhibit when exposed to high temperatures: thermoplastic and ther-

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moset polymers. When compared to thermosets, thermoplastic polymers are not as successful in terms of self-healing [2]. However, studies have been conducted in this area with promising results.

Thermoplastic Polymers Molecular Interdiffusion When two pieces of the same polymer are brought in contact with each other at temperatures above the material’s glass transition temperature, molecules diffuse across the interface and heals the “crack” which increases the mechanical strength of the polymer interface [11]. Figure 7 illustrates the diffusion process.

Fig. 7 Illustration of diffusion process at polymer/polymer interface (adapted from Zhang et al. [11])

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Fig. 8 Illustration showing dynamic de-cross-linking (adapted from Yuan et al. [13])

Photo-induced Healing In 2014, Sanjib Banarjee et al. showed the first instance of sunlight induced selfhealing material [12]. Low-molecular weight coumarin functional triarm star Polyisobutylene (PIB) was manufactured by a single step SN2 reaction of bromoallyl functional triarm star PIB with 4-methylumbelliferone or umbelliferone in the presence of sodium hydride [12]. When exposed to UVA, reversible photodimerization occurred in the coumarin moieties which formed cross-linked elastomeric films exhibiting self-healing behavior. The self-healing process was monitored by using atomic force microscopy after mechanical cuts were introduced in the films. Once irradiated, it was noted that the cuts healed up to 86% as the depth of the cuts decreased [12].

Self-healing via Reversible Bond Formation Chan’e Yuan et al. introduced a novel self-healing strategy that allowed dynamic cross-linking/de-cross-linking upon heating in cross-linked polystyrene without compromising its integrity [13]. Upon heating, covalent bond fission and radical recombination occur simultaneously among alkoxyamine moieties in the polystyrene chains and thus cracks are connected reconnected repeatedly, without losing integrity and load bearing ability of the material even above Tg [13]. This is shown in Fig. 8.

Thermoset Polymers Hollow Fiber Approach The basis of this approach is to fill brittle-walled vessels with polymerizable medium, which should be fluid at healing temperatures [14]. When the fluid medium flows to the area of damage, it polymerizes which leads to crack elimination. Three types

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Fig. 9 Schematic diagram of repair concept for polymer matrix composites using pre-embedded hollow tubes (adapted from Yuan et al. [14])

of healing systems have been developed. (i) Single-part adhesive. Hollow pipettes were filled with only one kind of resin such as epoxy particles (that flows once heated and then cured by the residual hardener) or cyanoacrylate (that hardens under the induction of air). (ii) Two-part epoxy. Epoxy and its curing agent are filled into neighboring hollow tubes, respectively. (iii) Two-part adhesive. One component is incorporated into hollow tubes and the other in microcapsules [14]. This concept is shown in Fig. 9. Motuku et al. investigated the hollow pipe reinforced healing systems in 1999 and found that species of healing agent, characteristic parameters of the hollow pipes (amount, type of tubing materials and spatial distribution), composites panel thickness, and impact energy level were found to be critical to the healing efficiency [15]. More recently, Trask et al. investigated the placement of self-healing HGF plies within both glass fibre/epoxy and carbon fibre/epoxy laminates for crack mitigation and mechanical strength restoration. It was found that even though there is an initial reduction in strength when the hollow fibers are incorporated into the material, there is an increase in damage tolerance. The material healed to 87% of its original strength [16].

Microcapsulation Approach This approach involves incorporating healing agents encased in microcapsules dispersed evenly throughout the polymer matrix along with a catalyst. When damage occurs, the microcapsules rupture and leak its contents into the area of the propagating crack and this leads to a chemical reaction between the healing agent and the catalyst leading to eventual healing of the material. Compared to other techniques, this approach has the most potential for commercialization as it does not require

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Fig. 10 Microencapsulated healing agent is embedded in a structural composite matrix containing a catalyst capable of polymerizing the healing agent (adapted from White et al. [18])

the change in molecular structure of the polymers to impart self-healing abilities and healing-agent-loaded microcapsules can be easily incorporated into the polymer matrix using existing blending techniques [17]. This is shown in Fig. 10. Currently, there are multiple self-healing systems that have been investigated and five different types have been proven to be most efficient. The single capsule system consists of only one type of healing agent encased in the microcapsules. Once the contents of the capsules are released into the site of damage, they react with the functional groups of the matrix activated by an outside agency such as moisture in the environment, light, heat, etc. to induce crack closure [17]. The capsule/dispersed catalyst system is made of microcapsules filled with monomers and a catalyst dispersed in the matrix. When the capsules break due to sustained damage, the monomers are released and polymerized with the help of the catalyst thus sealing the crack [17]. The phase-separated droplet/capsule system consists of at least one healing component that undergoes phase separation while the other component is encapsulated. The two components react with each other when they come into contact once rupture of the capsules occurs [17]. The double-capsule system employs at least two or more reactive healing agents that require compartmentalization [17]. Lastly, the all-in-one microcapsule system consists of the healing agent and the catalyst encapsulated in the same capsule but separated by layers or are encapsulated in smaller separate spheres and stored in a larger one. Once the capsules are ruptured, they come into contact thus generating the required reaction [17].

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Metals Most recent studies concerning self-healing have focused on polymers as they are easier to handle but metallic self-healing is an area of significant practical importance. Metallic atoms are very strongly bonded and have low diffusion rates, which make them difficult to work with. Several strategies have been suggested for metallic selfhealing as summarized below.

High-T Precipitation The original structure of the material of interest must contain a supersaturated amount of solute atoms, which can be attained by traditional metallurgical treatment if the two materials are compatible. When quenched, a metastable, supersaturated solid solution is formed that precipitates into secondary phases when aged further. The precipitation has to occur in localized regions where nano-voids are present rather than overall spontaneous precipitation [19]. A nano-void creates a grain boundary during damage, which results in a nucleation site for the precipitation process to occur as it attracts solute molecules to the site. High temperatures are required throughout the process to enable the lattice diffusion of the solutes towards the nano-voids [19]. Zwaag and group investigated this concept using high purity model systems of FeCu + BN and showed that the addition of B and N significantly accelerated Cu precipitation in a Fe–Cu alloy and most open volume defects (nano-voids) can be closed using this method [20].

Low-T Precipitation Compared to high-t precipitation, this procedure can be conducted at lower temperatures. It is very similar to the previous procedure as it requires the concept of supersaturated solute atoms but what differentiates is that the solute atoms tend to segregate to dislocation cores that can be considered as mobile even at low temperatures. Once nano-voids are formed due to stress, the solute atoms are attracted to these sites and eventually precipitate within the nano-voids to void closure. Al alloys supersaturated with Cu solute display self-healing properties that employ low-T precipitation; however, the material needs to be underaged to ensure that enough Cu atoms are left in solution to perform the process [19].

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NanoSMA-Dispersoids This concept is still in the developmental stage as it has been proposed fairly recently and the self-healing ability is yet to be confirmed. The host matrix needs to be embedded with shape memory alloy (SMA) nanoparticles that are stabilized in their austenite (high-temperature) phase. When damage occurs in the form of nano-voids or dislocations, the nanoparticles are activated. It is hypothesized that the stress field from the nano-voids trigger phase transformation of the SMA nanoparticles from austenite to martensite phase. This induces significant shape change in the particles that prompt local strain fields in the host matrix and leads to eventual crack closure [19]. Rajak and group have carried out simulations on precracked Al2 O3 embedded with SiC/SiO2 nanoparticles under mode 1 strain at 1426 °C. During crack propagation, nanoparticles in the vicinity of the damage create nanochannels that allow silica to flow towards the crack and impede its growth [21].

SMA-Clamp and Melt This SMA-clamp and melt concept is the most investigated macro length-scale concept. The structural features of interest on this scale are in the mm regime. The composite microstructure is made of SMA reinforced wires embedded in a solder matrix material. The solder acts as the “glue” and thus it should have a considerably lower melting point than the SMA wires. When the material comes under stress that exceeds the limits of the solder material, damage occurs to the matrix. The SMA wires transform to the martensite phase, instead as they have higher tensile strength than the matrix. The composite requires heating to temperatures above that required for the austenite transformation to trigger self-healing. This transition causes compressive stresses that contract the sample thus bringing the crack interfaces together [19].

Solder Tubes/Capsules The solder tubes/capsule concept incorporates solder material encapsulated inside ceramic capsules or ceramic tubes inside a host matrix with a higher melting temperature and significantly larger ultimate tensile strength. When damage occurs to the host matrix and a crack is formed, the solder is activated by increasing the temperature to its melting point which causes it to “wet” the crack surfaces fill it in based on capillary pressure and surface tension. Once the temperature is lowered, the solder material solidifies, closing the crack [19].

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Coating Agent The principle of self-healing coatings is the utilization of a coating that autonomically repair and prevent corrosion of the underlying substrate. To trigger the self-healing mechanism, heat treatment is required after which the crack is filled with the solder. It is similar to the solder tubes/capsules concept; however, the advantage is that no voids occur after the healing process [22].

Electro Healing The electro healing concept is different from other macro length-scale concepts as it requires no composite matrix or solder material. Instead, the damaged sample is immersed into an electrolyte solution and a voltage is applied. This induces and electrochemical reaction which causes the material to deposit inside the crack, eventually closing it [19]. The various phenomenon of self-healing in metals is summarized in Fig. 11 and Table 1.

Fig. 11 Schematic overview of the proposed/investigated self-healing concepts in metals (adapted from Grabowski and Tasan [19])

Nano

Management

Prevention

Solid-state healing

Solute precipitation

Creep resistance

Yes, at service temperature

Length scale

Nano scale damage

Macro scale damage

Type

Phase transition involved

Target property

Autonomous

High T

Precipitation

Yes

Fatigue resistance

Solute precipitation

Solid-state healing

Prevention

Management

Nano

Low T

Yes

No

All properties retained

Austenite ↔ martensite of wires

Austenite ↔ martensite of nano particles fatigue resistance

Liquid assisted

Management

Invisible

Macro

SMA-clamp and melt

Solid-state healing

Prevention

Management

Nano

NanoSMA dispersoids

No

?

Solidification of solder

Liquid assisted

Management

Invisible

Macro

Solder tubes/capsules

Table 1 Key features of the different self-healing concepts in metals (adapted from Grabowski and Tasan [19])

No

Fatigue resistance

Solidification of coating

Liquid assisted

Management

Invisible

Macro

Coating agent

No

Strength

Chemical reaction

Electrolyte assisted

Management

Invisible

Macro

Electro healing

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Ceramics Ceramics have strong and directional chemical bonds, which give their atoms limited mobility, and thus self-healing behavior is difficult to achieve. However, oxidative reactions at high temperatures can lead to reaction by products that can be used to fill small cracks. Ceramics usually procure most damage on their surface due to crash, fatigue, thermal shock, and corrosion during their service time as they are mostly used for structural use. Thus, self-healing of surface cracks in ceramics is an important issue that needs to be addressed to ensure the structural integrity of ceramic components [2]. In 2017, Toshio Oda and group demonstrated a novel approach to self-healing in ceramics [23]. They modeled their investigation on bone healing and followed the addition of a small amount of an activator, typically doped MnO localized on the fracture path, and found that it accelerated the healing time by more than 6000 times and also lowered the required reaction temperature. The activator filled the gap rapidly by generating mobile super cooled melts which allowed efficient oxygen delivery to the healing agent which further promoted crystallization of the melt and formed a mechanically strong healing oxide. The authors claim that this strategy can be used for the development of lightweight, self-healing ceramics that can be used on turbine blades in aircraft engines [23].

Self-healing Concrete and Asphalt Concrete technically falls under the ceramics category; however, it has vast potential for self-healing and thus it deserves its own section. Most of the damage that is sustained by cement is dominated by cracks. Self-healing in this context can be arranged into several broad categories: chemical encapsulation, bacterial encapsulation, mineral admixtures, chemical in glass tubing, and intrinsic self-healing with self-controlled tight crack width [2]. All these methods are summarized in Fig. 12. All these approaches have been successfully demonstrated in laboratory conditions; however, under natural conditions, the cement is introduced to uncontrolled limitations and thus further investigation is still required before a successful selfhealing cement can be developed.

Future Directions The development of self-healing materials has been mostly based on trying to emulate healing processes found in nature. Much advancement has been made in this area; however, there is still a long way to go if we are to replicate even the simplest biological healing mechanisms. Currently, the approaches that are available are single

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Fig. 12 a Chemical encapsulation self-healing approach, b bacteria additive self-healing approach, c mineral admixtures self-healing approach, d glass tubing self-healing approach, e self-controlled tight crack width self-healing approach (adapted from Li and Herbert [24])

step mechanisms [17]. Biological processes are much more complex, with multiple stages that occur simultaneously and the repaired site has properties that are similar to the original material. Wound healing in humans involves a multi-mechanistic approach in which the inflammatory response that is initiated by damage occurs at the same time as regeneration of the damaged material. In commercial synthetic routes, healing is still based on bridging or wedging as the only mechanism despite many advancements being made in other areas such as crack surface sliding and zone shielding [17]. Encapsulation and dispersion of healing agents within a matrix is very promising; however, the repaired material has inferior properties compared to the original matrix. Broader healing mechanisms need to be investigated and further limitations must be identified to achieve improved results. The biggest challenge to overcome is healing large cracks as current approaches can only manage cracks with small widths. As crack width increases, the efficiency of self-healing decreases significantly [25]. Further investigations need to be conducted in natural environments in which conditions cannot be controlled to obtain results that are more realistic. Despite

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these challenges, the trajectory of current research is opening doors to exciting possibilities in the future and the next-generation materials, with biomimetic healing abilities, will surely surpass the available technology. Conflict of Interest The authors have no conflict of interest to declare.

References 1. D’Elia E (2015) Self-healing organic/inorganic composites. Imperial College, London 2. Sobczak JJ, Drenchev L (2014) Self-healing materials as biomimetic smart structures. Foundry Research Institute 3. Garcia SJ (2014) Effect of polymer architecture on the intrinsic self-healing character of polymers. Eur Polym J (Elsevier) 53:118–125 4. Blaiszik BJ et al (2010) Self-healing polymers and composites. Annu Rev Mater Res 40(1):179–211 5. Fantner GE et al (2005) Sacrificial bonds and hidden length dissipate energy as mineralized fibrils separate during bone fracture. Nat Mater 4(8):612–6 6. Havaldar R, Pilli SC, Putti BB (2014) Insights into the effects of tensile and compressive loadings on human femur bone. Adv Biomed Res 3:101 7. Launey ME, Buehler MJ, Ritchie RO (2010) On the mechanistic origins of toughness in bone. Annu Rev Mater Res 40(1):25–53 8. Jackson AP, Vincent JFV, Turner RM (1988) The mechanical design of Nacre. Proc R Soc Lond Ser B Biol Sci 234(1277):415–440 9. Tran JP et al (2016) Bimaterial 3D printing and numerical analysis of bio-inspired composite structures under in-plane and transverse loadings, vol 108 10. Nosonovsky M, Rohatgi PK (2012) Thermodynamic principles of self-healing metallic materials. In: Biomimetics in materials science: self-healing, self-lubricating, and self-cleaning materials. Springer Series in Materials Science 11. Zhang H, Lamnawar K, Maazouz A (2012) Rheological modeling of the diffusion process and the interphase of symmetrical bilayer based on PVDF and PMMA with varying molecular weights, vol 51, pp 691–711 12. Banerjee S et al (2015) Photoinduced smart, self-healing polymer sealant for photovoltaics. ACS Appl Mater Interfaces 7(3):2064–2072 13. Yuan C et al (2011) Self-healing of polymers via synchronous covalent bond fission/radical recombination. Chem Mater 23(22):5076–5081 14. Yuan YC, Yin T, Rong MZ, Zhang MQ (2008) Self healing in polymers and polymer composites. Concepts, realization and outlook: a review. eXPRESS Polym Lett 2(4):238–250 15. Motuku M, Vaidya UK, Janowski CM (1999) Parametric studies on self-repairing approaches for resin infused composites subjected to low velocity impact. Smart Mater Struct 8:623–638 16. Trask RS, Williams GJ, Bond IP (2007) Bioinspired self-healing of advanced composite structures using hollow glass fibres. J R Soc Interface 4:363–371 17. Zhu DY, Rong MZ, Zhang MQ (2015) Self-healing polymeric materials based on microencapsulated healing agents: from design to preparation. Prog Polym Sci 49–50:175–220 18. White SR et al (2001) Autonomic healing of polymer composites. Nature 409:794 19. Grabowski B, Tasan CC (2016) Self-healing metals. In: Self-healing materials. Springer 20. Zwaag S (ed) (2007) Self-healing materials: an alternative approach to 20 centuries of materials science. Springer 21. Rajak P et al (2018) Faceting, grain growth, and crack healing in alumina. ACS Nano 12(9):9005–9010

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22. Yang Z et al (2011) The self-healing composite anticorrosion coating. Phys Procedia 18:216–221 23. Osada T et al (2017) A novel design approach for self-crack-healing structural ceramics with 3D networks of healing activator. Sci Rep 7:17853 24. Li VC, Herbert E (2012) Robust self-healing concrete for substainable infrastructure. J Adv Concr Technol 10:207–218 25. Meharie MG, Kaluli JW, Abiero-Gariy Z, Kumar ND (2017) Factors affecting the self-healing efficiency of cracked concrete structures. Am J Appl Sci Res 3(6):80–86

Part XLI

Solidification Processing of Light Metals and Alloys: An MPMD Symposium in Honor of David StJohn

Revealing the Heterogeneous Nucleation and Growth Behaviour of Grains in Inoculated Aluminium Alloys During Solidification Yijiang Xu, Daniele Casari, Ragnvald H. Mathiesen and Yanjun Li

Abstract An in situ X-ray radiographic study on the grain nucleation and grain growth of inoculated Al–10Cu and Al–20Cu alloys during isothermal melt solidification and directional solidification conditions with constant cooling rates has been carried out. The influence of additional level of inoculation particles, cooling rates, and temperature gradient on the nucleation rate and growth kinetics of grains have been quantitatively studied. The deterministic nature of the heterogeneous nucleation of aluminium grain on inoculant particles is revealed. Numerical microstructure models have been developed to simulate the nucleation and growth behavior of aluminum grains and a good agreement between the experimental results and simulation results have been achieved. Keywords Heterogeneous nucleation · Grain growth · Solidification · Aluminium alloy

Y. Xu · Y. Li (B) Department of Materials Science and Engineering, Norwegian University of Science and Technology (NTNU), 7491 Trondheim, Norway e-mail: [email protected] D. Casari · R. H. Mathiesen Department of Physics, Norwegian University of Science and Technology (NTNU), 7491 Trondheim, Norway © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_154

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Introduction Grain refinement by inoculation is a common practice during the casting and solidification of Al and Al alloys. As summarized in several review papers [1–4], the grain refinement mechanisms have been much understood through extensive studies in the past decades. Lots of experimental works have been made to study grain refinement behaviour in inoculated Al alloys, such as thermal analysis [5, 6], post-solidification characterization of cast samples, in situ X-ray diffraction [7–9] and in situ X-radiography study [10–13]. However, comprehensive in situ studies on the heterogeneous nucleation and grain growth under dedicated solidification conditions are still needed to reach an in-depth understanding of the nucleation kinetics under different solidification conditions (cooling rate, temperature gradient) and different solute/grain refiner addition levels, the nucleation ceasing mechanisms, and the grain growth kinetics. Besides, many numerical and analytical grain size prediction models have been developed [14–28] to better understand the nucleation and grain growth behaviour. However, most of the grain size prediction models [14, 16, 17, 24, 26, 29], are based on the assumption of spherical/globular grain growth kinetics. Thus, a more sophisticated model including globular to dendritic transition (GDT) and dendritic growth kinetics needs to be developed, by which the influence of grain morphology transition on the nucleation kinetics and predicted grain size can be quantitatively investigated. Furthermore, most of the published models intended for directional solidification [20, 24–26], are still based on a local isothermal melt solidification assumption. Therefore, new models with a rigorous treatment of grain nucleation on inoculant particles under the temperature gradient effect during is still demanded. In the present work, an integrated study by in situ X-radiography and numerical modelling will be carried out to reveal the heterogeneous nucleation and growth behaviour of grains during solidification of inoculated Al alloys.

Experimental The materials used in the experimental study are Al–10Cu and Al–20Cu (wt%) alloys prepared by melting 5N (99.999 wt%) purity aluminum and 4N (99.99 wt%) purity copper in a clay graphite crucible using a Nabertherm melting furnace. After complete melting and mixing of the raw metals, different levels of commercial Al–5Ti–1B (wt%) master alloy was added to the melt and finally cast into a copper mold. Thin plate samples with dimension of 5 × 50 × 0.2 ± 0.01 mm (X × Y × Z) are prepared from the cast ingots by cutting, grinding and polishing. The microfocus X-radiography setup applied in the present study, including the XRMON Gradient Furnace, CCD camera and X-ray source with a Mo transmission target, has been described in detail in Ref. [30–32]. The sample was aligned in a configuration where the broad surface (X-Y plane) of the sheet-like sample is

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perpendicular to the gravity (Z || g), by which the melt convection and grain motion were greatly reduced. The Bridgman furnace was operated in the so-called nearisothermal mode [12, 28], and also in standard directional solidification mode with imposed temperature gradients G along the sample length direction (G || Y). Constant cooling rates T˙ , in the range of 0.025–1.0 K/s, were applied by a controlled power down technique at two heater elements located at two sides of the field of view (FOV) along the sample length direction (Y direction). The images were recorded in situ at a frame rate of 2 Hz.

Model The numerical model used in this work is an extension to a previous grain size prediction model [26], now modified to include the dendritic growth and effect of temperature gradient on nucleation. The heterogeneous nucleation of grains on inoculant particles is based on the free growth criterion [17]. A particle-size distribution is measured as input parameters. Figure 1 is a schematic drawing to show the nucleation of new grains around one single grain during isothermal melt solidification and directional solidification. The solute diffusion field and temperature gradient induced inhibited nucleation zone (INZ, the red region) are rigorously treated based on the local undercooling shown in Fig. 1a1, b1, and particles can only be active for nucleation in the active nucleation zone (ANZ, the green region). The detailed model description and mathematical equations could be found in Ref. [28, 33].

Results and Discussion Isothermal Melt Solidification The in situ X-radiography image sequences during isothermal melt solidification of 0.05 wt% Al–5Ti–1B inoculated Al–20Cu alloy at three constant cooling rates, 0.025, 0.1 and 0.5 K/s are shown in Fig. 2. Due to the difference in Cu concentration between the liquid and solid phases, the solid Al grains show a brighter contrast. The nucleation kinetics and growth of each individual grain could be tracked from the image sequences. Under all these cooling conditions, equiaxed grain structures were obtained. However, as the cooling rate increases, the total number of grains appearing in the FOV also increases. The evolution of the total number of primary α-Al grains in the FOV as a function of undercooling below the nucleation temperature (since the first grain is observed) has been extracted and plotted in Fig. 3. Three stages nucleation kinetics can be observed. When the relative undercooling is small (≤~1 K), the numbers of grains formed in the FOV under different cooling rates are nearly the same at the same

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Fig. 1 Schematic drawing showing the nucleation of new grains around one single grain during a isothermal melt solidification and b directional solidification. a1, b1 Liquidus temperature Tl (r ), melt temperature T (r ) and the corresponding local undercooling of the liquid T (r ) outside the grain envelope that along the Y direction. a2, b2 Two-dimensional illustration of the inhibited nucleation zone (INZ), active nucleation zone (ANZ), and the corresponding boundary

undercooling values. This is because the volume fraction of solid grains is very low at small undercooling, and therefore the solute rejection from the growing grains has little effect on the nucleation, and the initiation of new grains only depends on the number of available potent inoculant particles. It is a strong evidence to support the athermal nucleation theory and free growth model [17], that the nucleation of grains is just a function of undercooling and independent of cooling rate at the beginning of solidification. However, at higher undercoolings, the number of grains in the FOV for different cooling rate is different and the nucleation rate is smaller in the low cooling rate cases. This means that the solute diffusion zone around growing grains, in terms of solute suppressed nucleation zone or inhibited nucleation zone, has played an important role in preventing the nucleation of new grains in the zones [20, 24]. In this stage, there is a competition between grain growth, solute segregation and external cooling. Once the real cooling rate is less than the reduction rate of liquidus temperature of the remaining melt (caused by the enrichment of solute in the residual liquid metal) [26], the nucleation process stops and the number density reaches the maximum. That is stage three and all the space between the grains become nucleation-

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Fig. 2 Selected X-radiographic images from the in-situ isothermal melt solidification of 0.05 wt% Al–5Ti–1B inoculated Al–20Cu alloy under three different cooling rates: a 0.025 K/s, b 0.1 K/s and c 0.5 K/s

free zones. It can also be seen from Fig. 2 that the average distance between neighbour grains and the areas of nucleation-free zones decrease with increasing cooling rate. Figure 4a shows predicted grain size of Al–10Cu alloy solidified under 0.5 K/s cooling rate, as a function of additional level of refiner, in comparison to the experimentally determined grain sizes from in situ X-radiography. As can be seen, the model prediction with considering the dendritic morphology transition based on hemispherical tip growth have a good agreement with the experimental results, showing the high sensitivity of the model to the grain refiner addition level. As addition level increases, grain size decreases. Figure 4b shows experimental determined and predicted grain size as a function of cooling rate for 0.05 wt% Al–5Ti–1B inoculated Al–10Cu alloy. As can be seen, the predicted grain sizes show quantitatively good agreement with the measured values. It confirms that the present model has a good prediction capability. Moreover, the solute segregation stifling mechanism and solute suppressed nucleation zone treatment proposed in the model is validated.

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Fig. 3 The evolution of total number of grains in the FOV as a function of undercooling below nucleation temperature for the three different solidification cases shown in Fig. 2

Fig. 4 a Grain size of Al–10Cu alloy solidified under 0.5 K/s cooling rate, as a function of additional level of refiner. b Grain size as a function of cooling rate for 0.05 wt% Al–5Ti–1B inoculated Al–10Cu alloy

Directional Solidification Figure 5 shows the X-radiographic images recorded during solidification of 0.05 wt% Al–5Ti–1B inoculated Al–20Cu alloy at a cooling rate of 0.1 K/s, with three different temperature gradients (G  0, 5, 10 K/mm) along the Y direction. By comparing these three conditions, we can find that for the higher G condition, the total number of grains is less and the grains become more elongated, and the length

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Fig. 5 Selected X-radiographic images from in-situ solidification of inoculated Al–20Cu alloy at 0.1 K/s cooling rate under three different temperature gradients G: a G Y ∼  0, b G Y  5 K/mm and c G Y  10 K/mm

of different primary dendrite arms of each grain is very different. Besides, the grain nucleation front of the directional solidification cases, as labeled by the dashed lines in Fig. 5b, c, propagates gradually towards the hot side of the FOV. The new grains formed in the nucleation front have the influence of blocking the growth of previously formed grains behind them by solute diffusion field impingement. For those grains without new grains forming ahead of them, the grain growth is free to continue, and accordingly, these grains will develop into elongated morphologies, e.g., Grain A and Grain B in Fig. 5c. The corresponding evolution of total number of grains in the FOV as a function of solidification time is shown in Fig. 6. As can be seen, the number of grains in the FOV increases with the solidification time, i.e., decreasing melt temperature, until a maximum value is reached and then remains constant. During near-isothermal melt solidification, the evolution of grain number in the FOV is quite smooth. However, the curves of directional solidification, especially at G  10 K/mm, show a stepterrace character, indicating that the propagation of nucleation front has wave-like

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Fig. 6 Evolution of the total number of grains in the FOV as a function of solidification time in three solidification cases shown in Fig. 5

nature. This is consistent with the finding by Prasad et al. [13, 34]. In isothermal melt solidification, a maximum number of 49 equiaxed grains is achieved in the FOV within 24 s, corresponding to an average nucleation rate of 2.97 mm−3 s−1 . However, in directional solidification, it takes 94 s and 147.5 s to form only 30 and 18 grains in the whole FOV for G  5 and G  10 K/mm cases, respectively. The corresponding average nucleation rate is 0.46 mm−3 s−1 and 0.18 mm−3 s−1 , respectively, which are about one order of magnitude smaller than that in isothermal solidification at the same cooling rate. The results show that the nucleation rate and the final grain number is reduced significantly when a temperature gradient is applied, although the cooling rate during the solidification is the same. The grain growth in G  0 and G  10 K/mm at 0.1 K/s are also quantitively analyzed from the in situ X-radiographic image sequences. Each dendrite arm is labeled with a number 1 for grains A–F, as illustrated in Fig. 5. The primary dendrite arm length is measured from the nucleation center of grains, and the results are plotted in Fig. 7. As can be seen, during isothermal melt solidification, e.g., Grain E and F, the growth of equiaxed grains stops earlier due to the soft impingement of solute diffusion field with that of neighbor surrounding grains. However, at G  10 K/mm, directional solidification, grain growth along the temperature gradient direction towards the hot side of the melt could last a long time due to the delay or lack of nucleation of new grains in front of the dendrite tip owing to the temperature gradient effects. For grain A and B, the growth velocity is almost constant. Figure 8 shows the predicted and measured grain number density as a function of temperature gradient for the inoculated Al–20Cu alloy solidified at 0.1 K/s. As can be seen, the model reproduces the experimentally determined evolution trend of

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Fig. 7 Evolution of individual primary dendrite arm length over time for several grains selected from in-situ X-radiographic images (labeled in Fig. 5) in two solidification conditions: isothermal and directional solidification with G Y  10 K/mm (T˙  0.1 K/s)

Fig. 8 Predicted and measured grain number density of the 0.05 wt% Al–5Ti–1B inoculated Al–20Cu alloy as a function of temperature gradient for the same cooling rate of 0.1 K/s

grain number density in relation to the temperature gradient G. The good agreement between the model prediction and experimental results proves that the proposed modeling approach to treat the inhibited nucleation zone and active nucleation zone around growing equiaxed grains under temperature gradient effects is feasible.

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Conclusions The kinetics of heterogeneous nucleation and grain growth in isothermal melt and directional solidification of Al–Ti–B inoculated Al–Cu alloys has been studied by in situ X-radiography. The effect of cooling rate, additional level of grain refiner, and temperature gradient are quantitatively investigated and revealed. Furthermore, a new grain size prediction model considering GDT and the effect of G on nucleation has been applied to simulate the grain nucleation and growth and a good agreement between the predicted results and experimental results are achieved. The free growth concept and the solute segregation stifling mechanism for grain nucleation have been verified by the in situ experiments and numerical modelling. Acknowledgements The financial support by The Research Council of Norway and industrial partners, for the PRIMAL project (project number: 236675), is gratefully acknowledged.

References 1. Murty BS, Kori SA, Chakraborty M (2002) Grain refinement of aluminium and its alloys by heterogeneous nucleation and alloying. Int Mater Rev 47:3–29 2. Quested TE (2004) Understanding mechanisms of grain refinement of aluminium alloys by inoculation. Mater Sci Technol 20:1357–1369 3. Easton MA, Qian M, Prasad A, StJohn DH (2016) Recent advances in grain refinement of light metals and alloys. Curr Opin Solid State Mater Sci 20:13–24 4. Greer AL (2016) Overview: application of heterogeneous nucleation in grain-refining of metals. J Chem Phys 145:211704 5. Johnsson M, Backerud L, Sigworth G (1993) Study of the mechanism of grain refinement of aluminum after additions of Ti- and B-containing master alloys. Metall Trans A 24:481–491 6. Johnsson M (1995) Grain refinement of aluminium studied by use of a thermal analytical technique. Thermochim Acta 256:107–121 7. Iqbal N, van Dijk NH, Offerman SE, Moret MP, Katgerman L, Kearley GJ (2005) Real-time observation of grain nucleation and growth during solidification of aluminium alloys. Acta Mater 53:2875–2880 8. Iqbal N, van Dijk NH, Offerman SE, Geerlofs N, Moret MP, Katgerman L, Kearley GJ (2006) In situ investigation of the crystallization kinetics and the mechanism of grain refinement in aluminum alloys. Mater Sci Eng A 416:18–32 9. Iqbal N, van Dijk NH, Offerman SE, Moret MP, Katgerman L, Kearley GJ (2007) Nucleation kinetics during the solidification of aluminum alloys. J Non-Cryst Solids 353:3640–3643 10. Reinhart G, Mangelinck-Noël N, Nguyen-Thi H, Schenk T, Gastaldi J, Billia B, Pino P, Härtwig J, Baruchel J (2005) Investigation of columnar–equiaxed transition and equiaxed growth of aluminium based alloys by X-ray radiography. Mater Sci Eng A 413–414:384–388 11. Nguyen-Thi H, Reinhart G, Mangelinck-Noël N, Jung H, Billia B, Schenk T, Gastaldi J, Härtwig J, Baruchel J (2007) In-situ and real-time investigation of columnar-to-equiaxed transition in metallic alloy. Metall Mater Trans A 38:1458–1464 12. Murphy AG, Mirihanage WU, Browne DJ, Mathiesen RH (2015) Equiaxed dendritic solidification and grain refiner potency characterised through in situ X-radiography. Acta Mater 95:83–89 13. Prasad A, McDonald SD, Yasuda H, Nogita K, StJohn DH (2015) A real-time synchrotron X-ray study of primary phase nucleation and formation in hypoeutectic Al–Si alloys. J Cryst Growth 430:122–137

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14. Maxwell I, Hellawell A (1975) A simple model for grain refinement during solidification. Acta Metall 23:229–237 15. Thévoz P, Desbiolles JL, Rappaz M (1989) Modeling of equiaxed microstructure formation in casting. Metall Trans A 20:311–322 16. Desnain P, Fautrelle Y, Meyer JL, Riquet JP, Durand F (1990) Prediction of equiaxed grain density in multicomponent alloys, stirred electromagnetically. Acta Metall 38:1513–1523 17. Greer AL, Bunn AM, Tronche A, Evans PV, Bristow DJ (2000) Modelling of inoculation of metallic melts: application to grain refinement of aluminium by Al–Ti–B. Acta Mater 48:2823–2835 18. Easton MA, StJohn DH (2001) A model of grain refinement incorporating alloy constitution and potency of heterogeneous nucleant particles. Acta Mater 49:1867–1878 19. Greer AL, Quested TE, Spalding JE (2002) Modelling of grain refinement in directional solidification. In: Schneider WA (ed) Light metals 2002. Minerals, Metals & Materials Soc, Warrendale, pp 687–694 20. Quested TE, Greer AL (2005) Grain refinement of Al alloys: mechanisms determining as-cast grain size in directional solidification. Acta Mater 53:4643–4653 21. Böttger B, Eiken J, Apel M (2009) Phase-field simulation of microstructure formation in technical castings—a self-consistent homoenthalpic approach to the micro–macro problem. J Comput Phys 228:6784–6795 22. Qian M, Cao P, Easton MA, McDonald SD, StJohn DH (2010) An analytical model for constitutional supercooling-driven grain formation and grain size prediction. Acta Mater 58:3262–3270 23. Men H, Fan Z (2011) Effects of solute content on grain refinement in an isothermal melt. Acta Mater 59:2704–2712 24. Shu D, Sun B, Mi J, Grant PS (2011) A quantitative study of solute diffusion field effects on heterogeneous nucleation and the grain size of alloys. Acta Mater 59:2135–2144 25. StJohn DH, Qian M, Easton MA, Cao P (2011) The interdependence theory: the relationship between grain formation and nucleant selection. Acta Mater 59:4907–4921 26. Du Q, Li YJ (2014) An extension of the Kampmann-Wagner numerical model towards as-cast grain size prediction of multicomponent aluminum alloys. Acta Mater 71:380–389 27. Martorano M, Aguiar D, Arango J (2015) Multigrain and multiphase mathematical model for equiaxed solidification. Metall Mater Trans A 46:377–395 28. Xu Y, Casari D, Du Q, Mathiesen RH, Arnberg L, Li Y (2017) Heterogeneous nucleation and grain growth of inoculated aluminium alloys: an integrated study by in-situ X-radiography and numerical modelling. Acta Mater 140:224–239 29. Quested TE, Greer AL (2004) The effect of the size distribution of inoculant particles on as-cast grain size in aluminium alloys. Acta Mater 52:3859–3868 30. Murphy AG, Browne DJ, Mirihanage WU, Mathiesen RH (2013) Combined in situ X-ray radiographic observations and post-solidification metallographic characterisation of eutectic transformations in Al–Cu alloy systems. Acta Mater 61:4559–4571 31. Nguyen-Thi H, Reinhart G, Salloum-Abou-Jaoude G, Browne DJ, Murphy AG, Houltz Y, Li J, Voss D, Verga A, Mathiesen RH, Zimmermann G (2014) XRMON-GF experiments devoted to the in situ X-ray radiographic observation of growth process in microgravity conditions. Microgr Sci Technol 26:37–50 32. Rakete C, Baumbach C, Goldschmidt A, Samberg D, Schroer CG, Breede F, Stenzel C, Zimmermann G, Pickmann C, Houltz Y, Lockowandt C, Svenonius O, Wiklund P, Mathiesen RH (2011) Compact X-ray microradiograph for in situ imaging of solidification processes: bringing in situ X-ray micro-imaging from the synchrotron to the laboratory. Rev Sci Instrum 82:105108 33. Xu Y, Casari D, Mathiesen RH, Li Y (2018) Revealing the heterogeneous nucleation behavior of equiaxed grains of inoculated Al alloys during directional solidification. Acta Mater 149:312–325 34. Prasad A, Liotti E, McDonald SD, Nogita K, Yasuda H, Grant PS, StJohn DH (2015) Real-time synchrotron X-ray observations of equiaxed solidification of aluminium alloys and implications for modelling. IOP Conf Seri Mater Sci Eng 84:012014

Influence of Microstructure Evolution During Twin-Roll Casting on the Properties of Magnesium Sheets K. U. Kainer, G. Kurz, S. Pakulat and D. Letzig

Abstract Twin-roll casting of magnesium alloys is seen as a promising processing route to enable further development of advanced magnesium sheets for mass production. The reduction in the number of processing steps to final gauge leads to shorter production times and a decrease in production costs. The new production process for magnesium sheets allows the development of a new generation of magnesium alloys. The microstructure evolution during solidification between the TRC rolls is of great importance and can be influenced by alloy composition, addition of grain refiner or by variation of the process parameter. In the twin-roll casting process, liquid metal is pumped from furnace or cast over a pipe into a tundish. The melt is then dragged into the roll gap of a pair of counter rotating, internally cooled rolls. The metal solidifies upon contact with the cooled rolls is rolled to a strip. Such strips are used as feedstock material for rolling sheets to final gauge. This presentation will discuss the effect of the influencing variables on the quality and performance of Mg sheets on the example of AZ31 sheets. The influence of the strip properties and the rolling process parameters like temperature and degree of deformation on the microstructure, the texture and the mechanical properties of the strip is presented and discussed. Keywords Magnesium sheet · AZ31 · Twin-roll casting · Rolling · Process parameters

Introduction At present, magnesium alloys used in the automobile industry are mainly processed by die casting. This technology allows components with a complex geometry to be manufactured. However, the mechanical properties of die-cast materials often do K. U. Kainer (B) · G. Kurz · S. Pakulat · D. Letzig Helmholtz-Zentrum Geesthacht – Magnesium Innovation Centre, Max-Planck-Straße 1, 21502 Geesthacht, Germany e-mail: [email protected] © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_155

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not meet essential requirements with regard to endurance, strength, ductility, etc. A promising alternative for thin, large area parts, such as automotive body components is to utilize sheet material. Sheet metal formed parts are characterized by superior mechanical properties and high-quality surfaces without pores in comparison to die-cast components. Substitution of conventional sheet materials such as steel or aluminum by magnesium sheets could lead to significant weight savings. However, it will be necessary to produce sheet material with competitive properties in an economic production process. Twin-roll casting is an economic production process for the generation of fine-grained feedstock materials that can subsequently be warm rolled to thin sheets. This production process for thin strips combines solidification and rolling into one single production step. Thus, it saves a high number of rolling and annealing passes in comparison to the conventional rolling process. The twinroll casting technology is already well established for example in the aluminum industry. However, applications to magnesium alloys are in their infancy at present [1]. Worldwide, there are a small number of industrial or laboratory scale twin-roll casters installed at universities, research facilities and companies. Initial results from these activities on conventional wrought and cast alloys have shown promising sheet properties [2–6]. The development of wrought magnesium alloys and their introduction into industrial, structural applications are the main goal of the activities at the Magnesium Innovation Centre MagIC of the Helmholtz-Centre Geesthacht (HZG). The current focus of the research work is on alloy design and the development of processing technologies for semi-finished magnesium products. In the particular case of sheet materials, it has been recognized that the feedstock for the warm-rolling process needs to be in the form of thin bands as they are produced via twin-roll casting, if thin magnesium sheets are to become competitive industrial products. For this purpose, HZG has installed a twin-roll caster. The twin-roll casting line consists of a furnace line (StrikoWestofen) and a twin-roll caster (Novelis, Jumbo 3CM). Important features of the strip such as microstructure and texture are influenced by the position of the solidification front and therefore the resulting degree of deformation as a result of the rolling pass. The position of the solidification front is controlled mainly by the melt temperature and the strip speed.

Experimental Procedure To investigate the influence of the melt temperature on the strip properties, twin-roll casting trials were carried out at temperatures between 680 and 710 °C at a casting speed of 1.8 m/min. The setback and the rolling gap were constant. In all trials, the commercial magnesium alloy AZ31 was used and strips cast with a width of 350 mm. In order to see the influence of the strip properties on the rolling result, strips were chosen which are twin-roll cast at the upper (715 °C 3.7 m/min) and the lower (650 °C 1.8 m/min) process limit. The AZ31 strip twin-roll cast at 650 °C was used in the as-twin-roll cast condition and rolled at three different rolling temperatures as well

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Table 1 Process parameters of the rolling trials [7] Temperature/Degree of deformation p. pass (°C)

ϕ  0.1 (10 passes)

ϕ  0.2 (5 passes)

ϕ  0.3 (4 passes)

350

AZ31 (650 °C) AZ31 (715 °C)

AZ31 (650 °C)

AZ31 (650 °C)

400

AZ31 (650 °C) AZ31 (715 °C)

AZ31 (650 °C) AZ31 (715 °C)

AZ31 (650 °C) AZ31 (715 °C)

450

AZ31 (650 °C) AZ31 (715 °C)

AZ31 (650 °C)

AZ31 (650 °C)

as three different degrees of deformation per pass. In order to compare the impact of the different microstructures in the feedstock material, strips of AZ31 twin-roll cast at 715 °C were rolled at 400 °C and at degrees of deformation of ϕ  0.1, ϕ  0.2 and ϕ  0.3. Additionally, rolling trials were performed at the three temperatures with a degree of deformation of ϕ  0.1. The processing parameters during rolling are noted in Table 1 [7]. Prior to the rolling procedure, the strips were reheated for 30 min to the respective rolling temperature. Between the following rolling passes, the rolled samples were again reheated to the rolling temperature for 15 min. After the final rolling pass, the sheets were air cooled. The rolling speed was 10 m/min and a water-soluble oil-based lubricant was used. After twin-roll casting and rolling, the microstructure of the material was analyzed using optical microscopy. Standard metallographic sample preparation techniques were employed and an etchant based on picric acid was used to reveal grains and grain boundaries of the strips and sheets [8]. Texture measurements in the strips and sheets were performed on the sheet midplanes using a Panalytical X-ray diffractometer setup and CuKα radiation. Six pole figures were measured up to a tilt of 70° which allows recalculation of full pole figures using an open source software routine MTEX [9]. The (0001) and (10–10) pole figures are used in this work to present the texture of the strips and sheets at midplane. In order to see how the different rolling procedures influence the mechanical properties of the sheets, tensile tests were performed according to DIN EN 10002. All samples were prepared in the sheet rolling direction and in the transverse direction.

Results of the Twin-Roll Casting Experiments Figure 1 displays micrographs of strip cast at a constant rolling speed of 1.8 m/min and varied melt temperatures (710, 700, 690 and 680 °C). It can be observed that the strips cast at 710 and 700 °C reveal a columnar structure where grains grow from the surfaces towards the central region of the strip. The central region itself reveals equiaxed grains in a band with certain thickness. A remarkable feature is the strong segregation band containing a large amount of impurities at the centreline

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of the strip. Such segregation is a common effect in this type of processing. The large amount of columnar grains is an indicator that the solidification front was located closer to the kissing point of the rolls and thus, the degree of deformation was small during the rolling operation due to incomplete solidification. Figure 1 also displays (0001) and (10–10) pole figures of the strip which both exhibit low intensities. A very weak alignment of basal planes parallel to the strip plane still gives hints to active deformation mechanisms, which start to form of a typical sheet texture. Interestingly, components with alignment in RD and TD are found also with low fraction. Furthermore, Fig. 1 shows the results for a casting temperature of 690 °C. The columnar type microstructure is found to be less significant if compared to the strip cast at 700 °C and grains are more equiaxed. Nevertheless, the strip cast at 690 °C is characterized also by a segregation band with a large amount of impurities at the centreline. The higher proportion of equiaxed grains in the microstructure indicates that the solidification front moves slightly towards the tip exit and the degree of deformation is somewhat higher than in the strip cast at 700 °C. Figure 1 also illustrates the (0001) and (10–10) pole figures of the strip. The result is very similar to that strip cast at 700 °C, which indicated comparable features in the centreline. It is noted that this position in the through-thickness consideration most likely is the least part influenced by active deformation. During casting, the melt temperature was reduced to 680° C (Fig. 1). This microstructure differs from the microstructures of the other three strips. It consists of smaller equiaxed grains, which are elongated in the rolling direction in the nearsurface regions of the strip. This microstructure is typical for a deformed material. Still, some features of a strip cast microstructure, as shown above, are visible. The respective pole figures of this strip are also given in Fig. 1. Again, the main feature

Fig. 1 Microstructure and texture after twin-roll casting with different melt temperatures at a casting speed of 1.8 m/min and a thickness of 5 mm

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is a weak alignment of basal planes in the sheet plane accompanied by a component tilted 90° to the RD as well as to the TD. Interestingly, there is no clear difference in the texture if compared to the results at higher casting temperature, which is a counterintuitive finding with respect to the consideration about the amount of realized deformation during strip casting and concurrent rolling. However, the measurement of the texture at midplane position only gives respect to the condition in the finally solidified fraction of the microstructure, which can be identified as a separate region in Fig. 1. Haßlinger et al. [10] showed that the development of an alignment of basal planes in the sheet plane is the result of continuously increasing deformation with increasing number of passes. At low degree of deformation comparable weak textures with the alignment of basal planes in the sheet plane have been found. Thus, the resulting texture reveals that the amount of deformation realized in this central part of the strip does not vary significantly with the casting temperature. Figure 2 depicts the microstructures and textures of samples taken from strips obtained from two different twin-roll casting trials [7]. Both pictures show the longitudinal sections of the strips. The microstructure on the right in Fig. 2 (right side) exhibits a columnar structure growing from the upper and bottom surfaces of the strip towards the center and a region containing equiaxed grains at the center [7]. Again, a remarkable feature is the strong segregation band containing a large amount of impurities at the centerline of the strip. The formation of this segregation band can be avoided only by optimization of the process parameters: casting speed, roll gap and melt temperature, which influences the thermal gradient. The large number of dendritic grains in the upper and lower regions of the strip is an indicator for a low degree of deformation during the rolling operation. Figure 2 also shows recalculated (0001) and (10–10) pole figures of these strips. In case of this high-temperature casting at high speed basically, a random texture is revealed at midplane. The random texture corresponds to no deformation at all, i.e. a globular cast microstructure. The microstructure and texture of a twin-roll cast strip with a melt temperature of 650 °C and a thickness of 5 mm are shown on the left side of Fig. 2. This microstructure differs significantly from the microstructures of the strip shown on the right side [7]. It consists of smaller equiaxed grains, which are elongated in the rolling direction in the near-surface regions of the strip. This microstructure is typical for deformed material [7]. Both the (0001) and (10–10) pole figures of the strip cast presented on the right in Fig. 2 are characterized by a weak alignment of basal planes parallel to the strip plane. Such a weak alignment of basal planes corresponds with a low degree of deformation [7]. In general, these results demonstrate that the strip cast at 650 °C and low-speed experienced more deformation than the strip cast at 715 °C and high speed [7].

Results of the Rolling Trials Figures 3 and 4 show the microstructures of the sheets rolled out of these strips [11]. It can be seen clearly that the microstructure of the rolled sheets changes

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Fig. 2 Microstructure and midplane texture of the alloy AZ31 after twin-roll casting different process parameters [7]

Fig. 3 Microstructure of the alloy AZ31 twin-roll cast at 715 °C after rolling at 350 °C and degree of deformation 0.1 [11]

significantly [11]. The more or less pronounced columnar microstructure at the top and bottom area disappears as well as the segregation band in the center. All sheets show a homogeneous, finer grained microstructure than the twin-roll cast strip before rolling [11]. The sheets cast at 715 °C and rolled with a deformation degree of ϕ  0.1 depict cracks after the 3rd/4th rolling pass at all temperatures [11]. The 715 °C cast strips rolled with deformation degrees of ϕ  0.2 and ϕ  0.3 could be rolled down to final gauge at 400 °C [11]. In all microstructures of the strip cast at 715° C, very distinct shear bands could be observed. Especially, the sheets rolled at a degree of deformation 0.1 show shear bands after the third and the fourth pass which leads to cracks in the surface (Fig. 3) [11]. These shear bands is often in direct neighborhood to long undeformed grains coming from the columnar microstructure of the twin-roll cast strip (Fig. 2) [11]. If the degree of deformation increases, the number of undeformed grains decreases and

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the sheets can be rolled to final gauge [7]. But, in comparison to sheets rolled out of the strip twin-roll cast at 650 °C these sheets exhibit a significant higher number of shear bands [7]. The (0001) and (10–10) pole figures of the AZ31 sheets cast at 715 °C are shown in Fig. 4 [11]. In comparison of the texture of the strip material cast at 715 °C, a pronounced basal texture of all sheets can be observed [11]. Figure 5 displays the microstructures from longitudinal sections of the as-rolled sheets of alloy AZ31, TRCed at 650 °C [12]. The microstructure of all sheets is not fully recrystallized. With increasing temperature the amount and size of the not recrystallized grains decreases. The sheets rolled at 450 °C with the degree of deformation of ϕ  0.1 are nearly fully recrystallized. With increasing degree of deformation, the amount and size of the deformed grains are increasing again. In the sheets rolled at 350 °C and a degree of deformation of ϕ  0.3, shear bands are observed. Like in the strips, there are some impurities spread over the whole thickness of the sheets. The (0001) and (10–10) pole figures of all AZ31 sheets cast at 650 °C are shown in Fig. 6 [12]. In comparison of the texture of the strip material, a pronounced basal texture of all sheets can be observed. The (0001) pole figures of nearly all the sheets exhibits split peaks towards the rolling direction in most cases. An angular spread to the transverse direction appears preferentially at the lower rolling temperatures which are consistent with the high amount of unrecrystallized grains in the microstructure [13]. In case of the sheet rolled at 450 °C and a degree of deformation of ϕ  0.1, one a single intensity peak and a broader angular spread to the rolling direction corresponds to a further recrystallized microstructure. Figures 7 and 8 summarize the results of these tests in stress–strain diagrams [12]. The stress–strain curves show in tendency that the elongation at fracture increases with increasing temperature, whereas the yield strength and ultimate tensile strength

Fig. 4 Microstructure and mid-plane texture of the alloy AZ31 twin-roll cast at 715 °C after rolling [11]

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decreases (Fig. 7). This inverse behavior can be observed at increasing degree of deformation. This material behavior is understood by a higher work hardening of the material at higher deformation degrees. At higher rolling temperatures, recrystallization effects weaken the material hardening and lead to increased formability of the sheet material, i.e. the fracture strain in the context of this work. The higher number of shear bands in the final sheets twin-roll cast at 715 °C leads to substantial worse mechanical properties (Fig. 8) [11].

Summary The results indicate that the processing parameters in twin-roll casting have a significant influence on the microstructure and texture of the strips. The degree of deformation in the strip depends on the position of the solidification front. If the solidification front is near the kissing point of the rolls, this leads to the effect that the upper and lower regions of the strip are solidified, but the central region is not completely solidified and therefore weak. Deformation during the rolling step is therefore concentrated in the weak central region without any effect on the outer regions of the strip. It could be shown that a melt temperature of 650 °C leads to a more fine-grained microstructure and to typical but weak rolling texture. The reason for this is, the solidification front is located near the exit of the tip, the degree

Fig. 5 Microstructures of the AZ31 sheets rolled at final gauge from strips twin-roll cast at 650 °C [11]

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Fig. 6 Textures of the AZ31 sheets rolled at final gauge from strips twin-roll cast at 650 °C [11]

Fig. 7 Stress–strain diagrams of the AZ31 sheets rolled at final gauge from strips twin-roll cast at 650 °C at different rolling temperatures [12]

Fig. 8 Stress–strain diagrams of AZ31 sheets rolled at final gauge from strips twin-roll cast at 650 and 715 °C [11]

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of deformation in the strip is increased. Therefore, the strip is completely solidified as it enters the rolls and can be deformed more homogenously. Rolling trials to final gauge sheet demonstrate that feedstock which received sustainable deformation with a homogeneous microstructure can provide better properties in the final sheet than material with a more pronounced cast structure and a resulting inhomogeneous microstructure. The results also reveal that the rolling process have a significant influence on the resulting sheet properties. It could be shown, that higher rolling degrees and lower process temperature lead to higher strength but lower formability in the sheet material. In contrast, higher rolling temperatures improve the formability but weaken the strength of the sheets. Acknowledgements This research was supported by funding from the German Federal Ministry of Education and Research. The authors are grateful to Novelis and STRIKOWestofen for assistance in the twin-roll casting trials.

References 1. Basson F, Letzig D (2010) Aluminium twin roll casting transfers benefits to magnesium. Alum Int Today 19–21 2. StJohn DH (2007) Overview of current international magnesium research and recent CAST CRC developments. J Adv Mater Res 29(30):3–8 3. Parks SS, Bae GT, Lee JG, Kang DH, Shin KS, Kim NJ (2007) Microstructure and mechanical properties of twin-roll strip cast Mg alloys. Mater Sci Forum 539(543):119–126 4. Kawalla R, Oswald M, Schmidt C, Ullmann M, Vogt HP, Cuong ND (2008) Development of a strip-rolling technology for MG alloys based on the twin-roll-casting process. TMS Magnesium Technology, New Orleans, Louisiana, pp 177–182 5. Watari H, Haga T, Paisarn R, Koga N, Davey K (2007) Mechanical properties and metallurgical qualities of sheets manufactured by twin-roll casting. Key Eng Mater 345(346):165–168 6. Aljarrah M, Essadiqi E, Kang DH, Jung IH (2011) Solidification microstructure and mechanical properties of hot rolled and annealed Mg sheet produced through twin roll casting route. In: 5th international conference on light metals technology. Light Metals Technology V, Lueneburg, Germany, pp 331–334 7. Kurz G, Bohlen J, Letzig D, Kainer KU (2013) Influence of process parameters on twin roll cast strip of the alloy AZ31. Mater Sci Forum 765:205–209. https://doi.org/10.4028/www. scientific.net/MSF:765.205 8. Kree V, Bohlen J, Letzig D, Kainer KU (2004) The metallographical examination of magnesium alloys. Pract Metallogr 5:233–246 9. Bachmann F, Hielscher R, Schaeben H (2010) Texture analysis with MTEX—free and open source software toolbox. Solid State Phenom 160:63–68 10. Haßlinger U, Fuskova L, Hartig C, Günther R, Letzig D, Kainer KU, Bormann R (2009) Effects of ceramic inoculants and rare earth phases on microstructure and related mechanical properties of magnesium alloys. In: Kainer KU (ed) 8th international conference on magnesium alloys and their applications, pp 1260–1267 11. Kurz G, Pakulat S, Bohlen J, Letzig D (2017) Influence of process parameters on the sheet properties of AZ31. In: FIMPART 2017, Bordeaux 12. Kurz G, Pakulat S, Bohlen J, Letzig D (2015) Rolling twin roll cast magnesium strips with varied temperature and degree of deformation. Mater Today Proc 39–44 13. Victoria-Hernandez J, Yi S, Bohlen J, Kurz G, Letzig D (2014) The influence of the recrystallization mechanisms and grain growth on the texture of a hot rolled AZ31 sheet during subsequent isochronal annealing. J Alloy Compd 616:189–197

A History of the Global Light Metals Alliance Jennifer Jackman, Kumar Sadayappan and Mark Easton

Abstract The concept for the Light Metals Alliance arose from discussions in 2001 among Prof. David StJohn (then Research Program Leader of CAST, Australia), Dr. Jennifer Jackman (then Director General of CANMET Materials Technology Laboratory) and Dr. Helmut Kaufmann (then Director of LKR, Austria). The objective of the Alliance is to enable better use of light metals in a broad range of real-world applications through collaboration and knowledge exchange. This is accomplished by exchanges of personnel for research and technology transfer. Members also organize a biennial conference (Light Metals Technology—LMT) showcasing government/industry/academia collaborations. Since 2001, membership has grown from 3 to 11, and there have been numerous scientific exchanges yielding more than 50 co-authored publications and 8 LMT conferences. Research covers all aspects of the production of parts from light metal alloys and composites, and the assessment of their mechanical and corrosion performance. The presentation will highlight success factors and achievements over the 17-year history of the Alliance. Keywords International collaboration · Light metals · Magnesium · Aluminum · Titanium

J. Jackman (B) CanmetMATERIALS (Retired), 212 Carmel Road, Castleton, ON K0K 1M0, Canada e-mail: [email protected] K. Sadayappan CanmetMATERIALS, 183 Longwood Road South, Hamilton, ON L8P 0A5, Canada e-mail: [email protected] M. Easton RMIT University, GPO Box 2476, Melbourne, VIC 3001, Australia e-mail: [email protected] © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_156

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Origins and Concept of the Alliance The Global Light Metals Alliance is an informal, international network composed of organizations that receive public funding and are accountable for public good outcomes; but whose mandate requires significant collaboration with the private sector and a responsibility to transfer technology for the commercial use of light metals. The GLMA supports two types of activities, the bi-annual conference Light Metals Technologies (LMT), and collaborative research among members. It was formally instituted in 2004 when five research organizations signed a Memorandum of Understanding outlining the objective, management and activities of the Alliance. Initially focussing on aluminum and magnesium in automotive applications for vehicle weight reduction, the scope of Alliance research has expanded to include titanium, and to encompass a range of applications including medical, aerospace and energy. Furthermore, by 2018, membership had expanded to eleven organizations from ten countries. The origins of the Alliance date from 2001. In that year, Dr. Helmut Kaufmann (who was then the Director of the Austrian Research Centre LEICHTMETALLKOMPETENZZENTRUM (LKR)) met Prof. David StJohn (who was then Research Program Leader of the Cooperative Research Centre for Cast Metals Manufacturing (CAST), based at the University of Queensland in Australia). Subsequently, they visited CanmetMATERIALS (then under the direction of Dr. Jennifer Jackman) in Ottawa, Canada. There were (and are) some interesting similarities between the scientific environment in Canada, Austria and Australia. All three countries have advanced economies where metals are important, yet they have relatively small populations, globally speaking. The research organizations under the direction of Prof. StJohn, Dr. Kaufmann and Dr. Jackman have expertise in metal processing; yet none of them is so large as to fail to appreciate the benefits of collaboration with others in order to have impact in an increasingly competitive world. LKR, Canmet and CAST all have a mandate to do practical research benefitting industry and society in general, and work in collaborative arrangements with the private sector, universities and other research institutions. In principle, the organizations were competitors; indeed, they had some common clients among the global companies that manufacture auto parts and assemble vehicles. The founders of the Alliance also felt strongly that there was a lack of research conferences that emphasize practical applications. Furthermore, they saw numerous advantages to working together collaboratively. The concept of the Light Metals Alliance was born in those discussions and further developed through correspondence over the next few months. The desire to develop mechanisms to stimulate interactions among the three research organizations lead them to meet again at King’s College in Cambridge in 2002, during the 8th International Conference on Aluminium Alloys. A photo marks the occasion (Fig. 1). Essentially the Alliance would commit to two streams of activity to showcase and conduct research through collaboration; the first being to support joint research among employees of the member organizations; and the second was to organize a conference series that would focus on real-world applications of light metals and

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Fig. 1 David StJohn, Jennifer Jackman and Helmut Kaufmann in Cambridge (UK), 2002, marking the occasion of the decision to formally create the Light Metals Alliance

bring Alliance members (and others) together every 2 years for energetic discussions on issues relating to the implementation of light metals in industrial applications. From the beginning, the Alliance has been a “coalition of the willing”. There is no source of funding for their activities other than what can be committed from their own budgets, or from project-specific proposals to funding organizations. It has an almost non-existent governance structure, and is very informal. There are no annual reports or multi-year plans. The representatives agreed to try to meet annually in the early years, usually at international meetings that they were already attending. During these “executive” meetings, the defining characteristics of the Alliance, and the conference series it would create, were further refined. The Australians were keen to organize the first Light Metals Technology conference [1], which was held in Brisbane in 2003 (even before a Memorandum of Understanding to create the Alliance was signed). Two other organizations soon joined––these were the Worcester Polytechnic Institute (WPI) in Massachusetts, USA (under the directorship of Prof. Diran Apelian) and the Magnesium Innovation Centre (MagIC) in GKSS in Geesthacht, Germany (under the directorship of Prof. Karl Kainer). By the time the first Memorandum of Understanding was signed in 2004, the Alliance had grown to five organizations.

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Growth of the Light Metals Alliance The GLMA does not seek to promote itself or to attain new members. Growth is not a specific goal, although it is open to new members that share the objectives of current members. Newcomers have been added to the Alliance, often by having collaborations with existing members, attending LMT conferences and seeing a good fit between the modus operandi of the Alliance and their own organizations. Membership does not bring any particular rewards, other than the opportunity to talk to other members and the obligation to help run an international conference. By the time a second MOU was signed in 2013 (9 years after the original MOU, bureaucracy not being a strong point of the Alliance), three new members were added: • The Light Metals Development Network (LMDN), under the Council for Scientific and Industrial Research (CSIR), South Africa (represented by Dr. Sagren Govender in 2018); • The Brunel Centre for Advanced Solidification Technology (BCAST), host to the UK Engineering and Physical Sciences Research Council (EPSRC) centre for Liquid Metal Engineering (LiME), Brunel University, United Kingdom (under the directorship of Prof. Zhongyun Fan); and • National Engineering Research Centre of Light Alloy Net Forming (NERC-LAF) of Shanghai Jiaotong University (SJTU), China (under the directorship of Prof. Wenjiang Ding). Given the obviously international nature of the Alliance, members began to refer to it as the Global Light Metals Alliance by 2005, although this name has not been formalized in the MOU. The MOU signed in 2018 included an additional three new members to bring the current membership to eleven (Fig. 2): • Korea Institute of Materials Science (KIMS), Republic of Korea (represented by Dr. Bong Sun You); • School of Engineering at Jönköping University, Department of Materials and Manufacturing, Sweden (represented by Prof. Anders Jarfors); and • The Light Metals and Manufacturing Research Lab (LMMRL) at The Ohio State University, USA (under the directorship of Prof. Alan Luo)

GLMA Conferences: Light Metals Technology The bi-annual meeting serves to propagate recent information relevant to the GLMA mandate, and to feature scientific research that enables practical use of light metals in society. Some thought went into developing a format that would present recent, high-impact research and encourage lively discussion during meetings. Organizers sought plenary speakers that could provide insightful overviews of industry directions and needs. An informal atmosphere was sought with a limited number of parallel sessions.

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Fig. 2 Representatives of the member organizations of the Global Light Metals Alliance in Pittsburgh, USA, 2017

The first such meeting was held in Brisbane in 2003 [1], which had a focus on developing automotive products from R&D on Al and Mg. In St. Wolfgang, Austria LMT2005 [2] emphasized the value chain, most recent research and industrial relevance. Over the years, LMT meetings grew both in scope and in numbers. The third conference in St. Sauveur Canada [3] exhibited an expansion in medical and aerospace applications, the inclusion of Ti, and sustainability considerations including life cycle assessment. In 2007, the conference returned to Australia, which at that time had a strong national program on light metals. The proceedings from LMT2007 [4] were published for the first time in the Materials Science Forum series. There were papers on a wide array of topics, and for the first time there were about the same number of papers on Ti as on Mg and Al. The fifth conference in the series was held in Lüneberg, Germany in 2011 [5]. As with preceding conferences, it provided information on the most recent applications, and included energy production and computational materials science as specific themes. LMT2013 [6] was held in Old Windsor, UK. With 161 papers, it was the largest conference to date, including a full span of applications, end-use considerations and novel materials such as nanocomposites. In 2015, the conference was held in Port Elizabeth, South Africa [7]. Among other subjects, there was an emphasis on the development of new alloys through innovative processing, and correlation with mechanical properties and applications. LMT2017 [8] was held in Pittsburgh, USA, as a session within the larger conference Materials Science and Technology 2017. In 2019, it is anticipated that the LMT conference will be held in China, organized by NERC-LAF in Shanghai, China. LMT conferences have included a significant number of presentations from industrial scientists; material that has not been previously presented; and keynote plenary presentations that outline opportunities and challenges affecting the use of light metal

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products (such as current drivers in the automotive or aerospace sectors) or global issues (such as sustainability or global supply). Presentations by Alliance members were sought but not highlighted in any way compared to non-members. For several LMT conferences, travel support was provided by an industrial (or public sector) sponsor to a student from each Alliance member country. The conferences are used as a networking opportunity for Alliance members, to this end an informal dinner for employees of LMA member organizations is typically organized either before or after the conference, which has the purpose of encouraging LMA members to meet each other and stimulate more collaborations. There is a meeting of the LMA “executive group” to discuss topics of interest. In total, the Light Metals Alliance have organized 8 conferences publishing 684 papers with an emphasis on innovative, fresh material and practical applications; representing a considerable contribution to the world’s knowledge of the processing, application and sustainability of light metals.

Collaborations The GLMA encourages collaborations among its members: supporting professional development of members’ employees and/or students; and achieving specific research goals defined on a project-by-project basis. Each project or secondment is governed by its own separate bilateral agreement. Hosting organizations attempt to provide accommodation for visitors from member organizations. There are some challenges to collaboration. The intent is to conduct research leading to commercialization which naturally brings with it concerns regarding intellectual property and commercial know-how. Information sharing among members is purely voluntary. All of the Alliance members have industrial clients who are understandably keen to avoid the release of some relevant discoveries. Secondly, the countries are widely separated geographically. Travel among them is expensive. In some cases, language is a barrier. There is no specific source of funding for their activities; each project must be independently funded. Nevertheless, the Alliance has managed to support a number of useful collaborations, including: • Some 28 projects involving at least 25 visits among the members, from periods of 2 weeks up to 13 months. The research resulting from these projects has produced more than 50 publications co-authored by Alliance members. • There have been exchanges of samples, collaborations by correspondence and successful co-applications for international research grants (Austria/Australia and UK/Germany). • Members have recruited one another’s students and provided enriching career development opportunities for both junior and senior employees. • Interactions have enabled, in addition to enriched research excellence, an informal benchmarking among the international members. • A sampling of research topics includes:

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– A substantial body of work on grain refinement and new rheocasting techniques (see comments by Prof. M. Easton below); – bolt-load retention (BLR) of Mg alloys (important to the assembly of products utilizing Mg parts), in which experimental and finite element modelling work was performed to determine and describe the stress relaxation (creep) of Mg alloys, and demonstrate the validity of a washer-load-cell design to measure BLR [9, 10]; – wrought Mg alloy development, including deformation and texture evolution during uniaxial loading and the influence of texture on corrosion and tensile properties [11, 12]. Other work involved the development and testing of wrought alloys with advanced formability; – effect of pressure and other process parameters on the quality of squeeze cast aluminum alloys [13]; – several projects on processing [14] and properties of magnesium sheet [15], including formability of sheet alloys, fracture propagation and toughness of sheet materials. Included is the development of a new, indirect test method for measurement of crack-tip opening angle (CTOA) that is suitable for automation (the CTOA concept was applied to automotive sheet metal for the first time, and was shown to be related to hole-expansion performance and fracture propagation resistance) [16]; – a number of projects on the casting of components, including a significant collaboration on the development of a test mold and investigations into the causes of hot tearing [17], which utilized thermomechanical simulation, and examined process parameters and the role of grain refinement. This work won two best paper awards [18, 19]. Other research included casting technology for high-integrity light metal castings [20], and dies for the determination of the castability of high-temperature alloys [21]; – research on the effect of ultrasonic processing on grain refinement and the properties of Al nanocomposites [22].

Reflections from Prof. Mark Easton The Global Light Metals Alliance (GLMA) became my global network of people with similar interests on impactful light metals research. My initial experience with the GLMA was through the Light Metals Technology conference series, the first being in 2003 in Brisbane where I remember sitting on a boat at the conference dinner talking to the Canadian delegates (having been born in Canada myself, although having grown up in Australia). I have been to many of the subsequent LMT conferences and these have really provided the basis for my global network. In 2004, I spent about 7 months at the Leichtmetallkompetenzzentrum, Ranshofen (LKR) in Austria (6 months work and a month’s holiday in Europe). I worked on a few projects on the topic of grain refinement. It was a very productive time with the publication of two journal papers from the work which have both been very

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well cited [23, 24] as well as three conference papers out of the work. The success of this work combined the knowledge of grain refinement developed within CAST with particular processing technologies that were being investigated by LKR, in particular, New Rheo-Casting and a grain refining master alloy developed by metal matrix composite technology that they had developed. This time at LKR also enabled the casting of a number of billets on the LKR Mg billet casting facility of a ZM20 alloy grain refined by a CAST-developed Mg–Zr grain refiner, now marketed as Micro-Zir, which was used extensively by colleagues at CAST for future extrusion alloy studies. This led to a relatively long collaboration between the CAST Co-operative Research Centre and LKR involving two visits to Australia by LKR staff to work on grain refinement of Mg alloys and later on a third visit on the design of automotive structures. Also, a colleague from CAST ended up working with LKR for approximately 5 years and later one of the CAST Ph.D. students worked at LKR for 3 years. During that time LKR was awarded a grant to support a collaboration between CAST in Australia and LKR titled “Austria–Australia collaboration on wrought magnesium alloys with increased formability” worth e300k. Another CAST Ph.D. student has ended up working for AMAG, which is a company co-located with LKR and now employs many of the people I worked with whilst I was at LKR. Another important collaboration through the GLMA has been with GKSS/HZG in Germany. The main outcome of that collaboration was through assessment of the castability and performance of magnesium high-pressure die casting alloys where we exchanged alloys for assessment. This resulted in a series of papers, in particular two journal papers [25] and [21]. The first of these papers was awarded the 2017 TMS Extraction and Processing Division Technology Award which is a major recognition of the success of the GLMA. In this work, staff at HZG was able to do alloy assessment that was not available at CAST (compression creep testing above 200 °C), whilst CAST was able to cast alloys under similar conditions including in a specially designed castability die and to do a suite of property testing of samples cast under the same solidification conditions. It was only because of the collaboration that the suite of alloys tested and the extent of the testing could have been undertaken, which lead to the award. I also hosted two students from Shanghai Jiaotong (and there have been a number of others that have been hosted throughout the CAST network). Both spent 1 year in Australia and this led to five journal publications and a conference publication. There have also been significant interactions with the Brunel Centre for Alloy and Solidification Technology where a number of former CAST staff and students have worked as well. There has been use of their melt shear technology at BCAST to grain refine the difficult-to-grain refine, CAST-developed extrusion alloy, AM-EX1, in the production of extrusion billet. The melt shear technology has also been used at CSIRO in their casting processes. These examples provide an illustration of the tangible benefits that have been achieved through the GLMA, apart from the general very good relationships that have been generated over the years, which have always been a highlight of the GLMA

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Light Metals Technology conferences and the other meetings where members have met up.

Concluding Remarks As an informal, largely unfunded network of publicly funded research organizations working on Al, Mg and Ti (and their alloys and composites), the Global Light Metals Alliance has been able to stimulate and facilitate a considerable body of effective research through encouraging collaborative research among its members and organizing the biennial conference series Light Metals Technology. This has been accomplished largely through collegial interactions, sharing of physical resources and knowledge, supporting travel of employees and students and ensuring exposure to holistic overviews of the opportunities and challenges to implementing light metals in practical applications. The collaborative characteristics of the Alliance that was established by its founders (Prof. David StJohn being one of the principals) are still evident among the current eleven members, and stand in contrast to the competitive interaction that might otherwise be expected.

References 1. Dahle A (ed) (2003) Proceedings of the first international light metals technology conference. CRC for CAST Metals Engineering, Brisbane, Australia. ISBN 0-9751329-0-3 2. Kaufmann H (ed) (2005) Proceedings of the second international light metals technology conference. ARC Leichtmetallkompetenzzentrum Ranshofen GmbH (LKR), St. Wolfgang, Austria. ISBN-3-902092-03-3 3. Sahoo M, Sadayappan K (eds) (2007) Proceedings of the third international conference on light metals technology (LMT2007). Public Works and Government Services Canada, St. Sauveur, Canada. ISBN NR15-76/2007E 4. Dargusch MS, Keay SM (eds) (2009) Proceedings of the 4th international conference light metals technology 2009, vol 618–619. Trans Tech Publications Ltd. Switzerland, Materials Science Forum, Gold Coast, Australia. ISSN 0255-5476 5. Dieringa H, Hort N, Kainer KU (eds) (2011) Proceedings of the fifth international light metals technology conference (LMT2011), vol 690. Trans Tech Publications Ltd., Switzerland, Materials Science Forum, Lüneberg, Germany. ISSN 0255-5476 6. Stone I, McKay B, Fan Z (eds) (2013) Light metals technology 2013, vol 765. Trans Tech Publications Ltd., Switzerland, Materials Science Forum, Old Windsor, Great Britain 7. Chikwanda HK, Chikosha S (eds) (2015) Light metals technology 2015, vol 828–829. Trans Tech Publications Ltd., Switzerland, Materials Science Forum, Port Elizabeth, South Africa. ISSN print 0255-5476 8. (2017) Light metals technology 2017. Materials Science and Technology 2017 (MS&T17), Pittsburgh, USA, pp 12–177 9. Anopuo O, Shen G, Xu S, Hort N, Kainer KU (2009) Elevated temperature and varied load response of AS41 at bolted joint. In: Nyberg E et al (eds) TMS annual meeting & exhibition. Magnesium Technology, San Francisco, California, pp 509–514

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10. Xu S, Shen G, Williams G, Anopuo O, Hort N, Kainer K-U, Staron P (2012) Development of a standard test to evaluate bolt-load retention of magnesium alloys. In: Poole WJ, Kainer KU (eds) 9th international conference on magnesium alloys and their applications (ICMAA2012), Vancouver, Canada, pp 689–694 11. Yi S-B et al (2006) Deformation and texture evolution in AZ31 magnesium alloy during uniaxial loading. Acta Mater 54:549 12. Balzereit S (2012) Investigations of the influence of the crystallographic texture on the corrosion behaviour and the tensile properties of magnesium alloys. Diploma thesis, University of Technology, Hamburg 13. Fragner W, Sadayappan K (2007) Effect of processing on the structure and properties of squeeze cast Al-7Si-0.3 Mg alloy. In: Creapeau J, Campbell J, Tiryakioglu M (eds) 2nd international symposium on shape casting. The Metallurgical Society 14. Dieringa H et al (2015) Twin-roll casting after intensive melt shearing and subsequent rolling of an AM30 magnesium alloy with addition of CaO and SiC, vols 828–829. Material Science Forum, Port Elizabeth, South Africa, pp 35–40 15. Sotirov N et al (2007) Rolling of microalloyed magnesium sheets. In: Public works Canada, 2007. Light Metals Technology, St. Sauveur, Canada, pp 229–233 16. Xu S, Petri N, Tyson WR (2009) Evaluation of CTOA from load vs. load-line displacement for C(T) specimen. Eng Fract Mech 76:2126–2134 17. He Y et al (2013) Thermomechanical simulation and experimental characterization of hot tearing during solidification of aluminum alloys. Int J Cast Metals Res 26:72–81 18. Li S, Apelian D, Sadayappan K (2012) Hot tearing in cast al alloys-mechanisms and process controls. Int J Metalcast 6:51–58 19. Li S, Apelian D, Sadayappan K (2012) Hot tearing in cast aluminum alloys; mechanisms and process controls. In: 116th AFS casting congress, Columbus, USA 20. Kaufmann H, Uggowitzer PJ (2007) Metallurgy and processing of high-integrity light metal castings. Fachverlag Schiele & Schön GmbH, Berlin 21. Easton MA et al (2016) Evaluation of magnesium die-casting alloys for elevated temperature applications: castability. Adv Eng Mater 18:953–962 22. Wang G et al (2014) The role of ultrasonic treatment in refining the as-cast grain structure during the solidification of an Al-2Cu alloy. J Cryst Growth 480:119–124 23. Easton MA et al. Grain refinement of Mg-Al(-Mn) alloys by SiC additions. Scr Metall 55:379–382 24. Easton MA, Kaufmann H, Fragner W (2006) The effect of chemical grain refinement and low superheat pouring on the structure of NRC castings of aluminium alloy Al-7Si-0. Mater Sci Eng A 420:135–143 25. Zhu SM et al (2015) Evaluation of magnesium die-casting alloys for elevated temperature applications: microstructure, tensile properties, and creep resistance. Metall Mater Trans A 46:3543–3554

Analysis of the High-Purity Aluminum Purification Process Using Zone-Refining Technique Heli Wan, Baoqiang Xu, Jinyang Zhao, Bin Yang and Yongnian Dai

Abstract The article presents the results of an experimental study of the effect of impurity transport in a zone-refining system. In order to further improve the purity of aluminum, therefore, a sample with the purity (99.99%) aluminum material was used for zone refining, and evaluated its purity by glow discharge mass spectrometry (GDMS). In addition, the distribution of impurities after approximately 5 passes of zone refining was analyzed. In experiments, when the zone speed was increased from 5 to 30 mm/min at a zone width of 40 mm, and the zone refining were controlled 5 to 20 passes. The experiment result showed that impurities content was significantly affected by the difference in the number of zone passes. In addition, with the decrease of the moving speed of the melting zone, the content of impurities in the product was also gradually decreased by GDMS analysis. Keywords Zone refining · Zone speed · Zone pass · Impurities analysis · Aluminum

H. Wan · B. Xu (B) · J. Zhao · B. Yang · Y. Dai National Engineering Laboratory for Vacuum Metallurgy, Kunming University of Science and Technology, Kunming 650093, People’s Republic of China e-mail: [email protected] H. Wan · B. Xu · J. Zhao · B. Yang Research Center of Engineering on Aluminum Industry of Yunnan Province, Kunming 650093, People’s Republic of China H. Wan · B. Xu · J. Zhao · B. Yang · Y. Dai Yunnan Provincial Key Laboratory for Nonferrous Vacuum Metallurgy, Kunming University of Science and Technology, Kunming 650093, People’s Republic of China H. Wan · B. Xu · J. Zhao · B. Yang · Y. Dai Faculty of Metallurgical and Energy Engineering, Kunming University of Science and Technology, Kunming 650093, People’s Republic of China © The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6_157

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Introduction High-purity aluminum compared with aluminum has perfect ductility, conductivity, light reflectivity, and corrosion resistance, so it can be used for the preparation of aluminum film, cathode sputtering target materials and electronic components [1, 2]. At present, the purification of aluminum mainly adopts electrochemical method and physical method of segregation method, which includes three-layer liquid electrolysis method and organic solution electrolysis method. The physical methods include vacuum evaporation, fractional crystallization, directional solidification and regional melting [3–5]. Three-layer liquid aluminum electrolysis refining method is to purify the metal according to the electrode potential difference in the electrolyte [6, 7]. Physical segregation method, also known as coagulation purification method, is a purification process by using segregation phenomenon produced by the melt during solidification. The main methods are condensation, fractional crystallization, unidirectional solidification and regional melting [8–10]. Zone refining is an unconventional method of applying coupled melting and freezing processes for purifying metals, semiconductors, organic and inorganic chemicals. It was first proposed by Pfann [11, 12]. Zone refining of aluminum was first reported by Bratsberg et al. [13, 14]. Kino and co-workers discussed in detail about the relationship between zone-refining parameters and the purification efficiency from a 5N starting material [15, 16]. Hashimoto and co-workers reported super highly purified aluminum from a 6N starting material by the cropping method [17, 18]. The aim of the present work is to combine a theoretical and an experimental analysis in order to investigate the effect of experiment conditions on the efficiency of removal efficiency of impurities which can be used for planning a successful zone-refining procedure. Experimental work was carried out by using Al samples. Axial impurities profiles have been experimentally determined after a number of molten zone passes, and these profiles were compared with theoretical analyzes. The influence of such approach on the maximum purification achieved by the zonerefining technique, and the ultimate impurities content is also analyzed.

Experimental Raw materials used in this study included aluminum (>99.9% purity, provided by XINJIANG JOINWORLD) which is the composition of the aluminum as shown in Table 1, high pure argon gas (>99.99% purity). The samples, of length 300 mm, are then placed in a quartz boat inside a quartz tube and the system is evacuated and sealed. The quartz tube is loaded into a horizontal zone refiner equipped with one zone heaters, and accepting samples with a maximum length of 1000 mm. A

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Table 1 Composition of raw materials used in experiments Composition of elements (wt%) Al

Fe

Si

Cu

Zn

Ti

Mg

99.997

0.0007

0.0005

0.0013

0.0002

0.0001

0.0009

high-frequency resistance heating technique is used. The purification experiments are carried out under conditions of a constant 40 mm zone length and a zone travel rate of 5–30 mm/min. After 5, 10, 15 and 20 passes of the molten zone, the samples are examined for purity using GDMS.

Results and Discussion The Theoretical Analysis In zone melting process, the metal is prepared as a rod or ingot, and placed in a sealed quartz tube. The quartz tube is heated by an electromagnetic induction coil, and the metal bar or ingot will undergo a melting and solidification process. The part of the rod metal is heated through the electromagnetic induction coil, so that a small portion of the rod metal is melted in the coil wrapping part. Impurities migrate to the liquid–solid interface, as shown in Fig. 1. When the equilibrium distribution coefficient of impurity element k0 < l, the impurities move with the molten zone moving and finally migrate to the ingot metal head. Meanwhile, as for the equilibrium distribution coefficient of impurity element k0 > 1, the impurities migrate in the opposite direction, and finally concentrate to the tail of the sample. The sample is purified by zone melting and the diffusion of impurities is change as shown in Fig. 2. After multi-passes zone melting, the zone purification efficiency is especially distinctive [19, 20]. In the process of zone smelting, the zone-refining efficiency is affected by these factors, i.e., the effective distribution coefficient keff , zone melting speed of V and melting number N. The influence of keff is dominant during the purification process, which is adopted to describe impurity segregation, defined as follows [21]:

Fig. 1 Sketch of impurities segregation during purification process

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Fig. 2 Schematic diagram of the melting zone

Table 2 Concentrations and distribution coefficients of main elements in the raw material Element

Concentration, C/at ppm

Fe

7

0.03

Cu

13

0.14

Zn

2

0.75

Ti

1

2.1

keff 

Distribution coefficient, k0

k0 , (1 − k0 )e−(νδ/D) + k0

(1)

where k0 is the equilibrium distribution coefficient, ν is the molten zone traverse velocity, δ is the thickness of diffusion boundary and D is the impurity diffusion coefficient in the melt. When ν  D/δ, keff → 1, which means that the impurity distribution in the solid is uniform, and that the refining efficiency is poor. When ν  D/δ, keff → 0, which indicates that the segregation efficiency is dependent on the equilibrium distribution coefficient of a specific impurity. Analytical solutions of impurity distribution in the multi-pass zone-refining process are based on the above factors. The impurity content and distribution coefficient in the raw material are shown in Table 2 [22, 23]. Therefore, combining with the content and characteristics of impurities the current experimental materials, the zone smelting technology is adopted in this paper to study and analyze the removal of impurities (Fe, Cu, Zn and Ti).

Influence of the Melting Zone Speed Figure 3 indicates that the content of impurity Fe is gradually decreased with the decreasing of moving zone speed, and finally concentrated in the tail of the sample. When the melting zone speed is 5 mm/min, the concentration of Fe in the middle part of the sample is approximately 1.1 ppm, which meet the standard of 5 N alu-

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Fig. 3 Concentration distributions of impurities along the sample for zone refining of Al with different the melting zone speed: a Fe, b Cu, c Zn and d Ti

minum(YS/T 275-2008, Fe ≤ 2.5 ppm), and the concentration of the impurity iron is increased at the tail end. However, with molten zone speed reducing, the removal effect of the impurity copper is not obvious. Although with the slowing down of the movement speed of the melting zone, the purification efficiency is continuously improved, and the Cu is also constantly migrating towards the tail of sample. It shows that the width of melting zone and the number of melting zone pass are dominant factors which influenced the efficiency of purification. Therefore, the concentration of Cu is still as much as 5.3 ppm when the speed of melting zone is 5 mm/min, as shown in Fig. 3b. During a stable 10-pass zone-refining process, the concentration distribution of the impurity Zn at different locations in the sample with different speeds as shown in Fig. 3c. The results show that the Zn migrates to the tail of sample after several times of zone melting, and the Zn content in the first 40% part of the sample is lower than 0.88 ppm. The results also show that zone-refining efficiency significantly with the melting speed decreasing, the Zn concentration in the sample can be reduced to 0.32 ppm when the moving speed of the melting zone is 5 mm/min. In contrast, the purification effect is significant after multi-pass zone melting, and the Ti is concentrated in the top, and the Ti content is only 0.12 ppm in the tail sample as shown in Fig. 3d.

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One can see that purification efficiency and distribution position of impurities are different, which is significantly influenced by the equilibrium distribution coefficient (k0 ), Fe, Cu and Zn, which k0 are 0.03, 0.14 and 0.75, respectively. The concentration of impurities (Fe, Cu and Zn) are gradually increased along the ingot axial direction, however, the concentration of the Ti gradually decreased which k0 is 2.2. The experimental results are in good agreement with the reported by Masayoshi and Nakamura [24, 25].

Influence of Zone Pass During the experiment, the influence of zone melting pass on the efficiency of the purification is investigated in 5, 10, 15, 20 times, respectively. Figure 4 showed a comparison of experimental results with different zone pass and the concentration of impurities the zone refining of Al samples and now with Fe, Cu, Zn and Ti as the impurities, the zone travel rate is controlled at 5 mm/min. It can be seen that the impurities are concentrated at both ends of the sample, i.e., Fe, Cu, Zn of k0 < 1 are concentrated at the tail of the sample, respectively, as shown in Fig. 4a–c. However, Ti of k0 > 1 is concentrated at the top of the sample as shown in Fig. 4d. To improve the efficiency of the purification process, it is concentrated on impurities with k0 < 1 specifically for the Fe with k0  0.03 and Zn with k0  0.75 to obtain the optimized number of passes. As shown in Fig. 4a, c, the impurities concentration is analyzed by GDMS. The results shows that, after 15 passes, the purification process approaches the saturation value (CFe  1.35 ppm, CZn  0.47 ppm, CTi  0.13 ppm). For the impurity Cu, the efficiency of purification is again overestimated when the zone passes are only used as an optimization parameter in experiments. CCu is 5.09 ppm which content is still higher than the standard value (Cstandard  2.8 ppm), even for a 20 zone passes as shown in Fig. 4b. The results show that if the separation efficiency of CCu is improved, it is necessary to further increase the number of zone smelting and reduce the moving rate of melting zone. Figure 4d examined the experimental distribution of Ti during the purification of Al after multi-pass of the molten zone. Most of the experimental pointed that the impurity Ti is mainly concentrated in the top of the ingot. However, if zone refining after 15 passes is considered, the observed differences are not significant with increasing the number of zone melting pass.

Influence of Length of the Sample The width of melting zone 40 mm is used in the experiments, and Al ingot of length 260 mm is selected for melting and purification experiments which are mainly limited by the size of the graphite boat (L  300 mm). Figure 5a showed that the sample

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shape changed with the zone speed is 5 mm/min after 15 passes of zone melting. The solidified sample is cut into several pieces, including the top and tail portions shown in Fig. 5b. We carry out composition analysis with respect to the black points in the Fig. 5b. The chemical compositions of solute atoms are estimated by GDMS measurement. As mentioned in Sect. 3.3, the purification efficiency is improved when the molten passes are increased. Therefore, we analyze the effect of the length of the sample on the concentration profile and the results are displayed in Fig. 6. For any conditions examined in the present study, the impurities (Fe, Zn, Cu) content are decreased

Fig. 4 Concentration distribution of impurities along the sample for zone refining of Al with different passes: a Fe, b Cu, c Zn and d Ti

Fig. 5 Positions used for composition analysis

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Fig. 6 Concentration distribution of impurities along the sample for zone refining of Al (ν  5 mm/min, n  15 passes)

Table 3 The impurity content of high-purity aluminum

Impurity elements

The content of elements (ppm) Fe

Zn

Cu

Ti

Al-5N (YS/T 275-2008)

2.5

0.9

2.8

1.0

(0.3–0.4) sample

1.49

0.5

5.61

0.35

gradually along the direction of melting zone, but Ti content is the opposite. It is mainly affected by the equilibrium distribution coefficient k0 of the impurity itself, which are in good agreement with the simulation results reported [26]. It can be seen that experimental results after 15 zone passes, the contents of impurities (Fe, Zn, Ti) are low except Cu in the 16.6% part (0.3–0.4) of the sample, as shown in Table 3. In the purification process, since solute transfer takes much longer time to reach stable than heat transfer takes time. Therefore, the purification efficiency is further improved and the content of Cu impurities is reduced. Considering solid and liquid concentrations in equilibrium at the eutectic temperature, one can further see that boundary layer thickness is also an important factor affecting the zone purification efficiency. From theory of fluid mechanics, the boundary layer thickness is dependent on convections [27]. It is necessary to optimize the experimental conditions, i.e., length of the sample and change the speed of the melting zone, which improves segregation of the efficiency of Cu.

Conclusions In the paper, we carried out zone melting experiments with different experimental conditions, and the sample after melting was analyzed by GDMS. The results obtained in this study are as follows.

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(1) When the molten zone speed reduced from 30 to 5 mm/min, after 10 zone passes, the concentration of impurity elements of k0 > 1 elements (Fe, Cu, Zn) in the last-half portion of the refined sample, the concentration of impurity elements of k0 > 1 elements (Ti) in the first-half portion of the refined sample. (2) The effect of molten zone passes on the refinement became prominent when the zone speed was controlled at 5 mm/min. The results shows that, after 15 passes, the purification process approaches the saturation value (CFe  1.35 ppm, CZn  0.47 ppm, CTi  0.13 ppm, CCu  5.09 ppm). (3) The GDMS results of Cu was not in good agreement with the theoretical analysis even when the zone speed and/or zone pass varied. This finding indicates the usefulness of improving the separation efficiency of Cu in our experimental method. Acknowledgements This work was supported by National Science Foundation of China (No. 51734006), Science and Technological Talent Cultivation Plan of Yunnan Province, China (No. 2017HB009), the Cultivating Plan Program for the Leader in Science and Technology of Yunnan Province under Grant (No. 2014HA003) and the Program for Nonferrous Metals Vacuum Metallurgy Innovation Team of Ministry of Science and Technology under Grant (No. 2014RA4018).

References 1. Jia ZH, Ma MZ, Xu L (2015) Effect of segregation purification process on content of impurity elements in high purity aluminum. Adv Mater Res 1061–1062:71–75 2. Ezheiyan M, Sadeghi H (2017) Simulation for purification process of high pure germanium by zone refining method. J Cryst Growth 462:1–5 3. Kamavaram V, Mantha D, Reddy RG (2003) Electrorefining of aluminum alloy in ionic liquids at low temperatures. J Min Metal 39(1–2):278–281 4. Production of extreme purity aluminum (1978) Aluminum company of arncrica. US Patent 2039529A, 26 Dec 1978 5. Jin JZ, Ren ZM, Yu GL (1985) A device for purifying aluminum by directional solidification. C A Patent No 85201157U, 2 1985 6. Kondo M, Maeda H, Mizuguchi M (1990) The production of high-purity aluminum in Japan. JOM 42(11):36–37 7. Hashimoto E, Ueda Y, Kino T (1995) Purification of ultra-high purity aluminum. J Phys IV 5(C7):153–157 8. Yan FU, Zhang XM, Shun ZQ (2001) Purification of aluminium ingot by segregation method. J Northeastern Univ 22(2):136–139 9. Rambabu U, Munirathnam NR, Reddy RC (2005) Segregation behaviour of trace metal impurities during ultra high purification of gallium by zone refining. J Pure Appl Phys 43(10):783–786 10. Zhao RM, Zhang LY, Yang G (2016) Segregation behavior of impurity elements in segregation purification of aluminum. Light Alloy Fabric Tech 12:23–25 11. Pfann WG (1952) Principles of zone-melting. JOM 4(7):747–753 12. Pfann WG (1966) Zone melting, seconded. Wiley, New York 13. Bratsberg HG, Herbjornsen OH, Foss D (1963) Zone refining of aluminum. Rev Sci Instrum 34(7):777–778 14. Revel G (1964) Aluminium de haute purete obtenu par zone fondue. Comptes Rendus 259(22):4031–4033 15. Kino T, Kamigaki N, Yamasaki H (1976) Zone refining of aluminum. Trans JIM 17(10):645–648

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16. Kino T, Hashimoto E, Kamigaki N (1977) Study on the trace elements in zone-refined aluminum. Matsushita Trans JIM 18(4):305–312 17. Hashimoto E, Ueda Y (1994) Zone refining of high-purity aluminum. Mater Trans JIM 35:262–265 18. Hashimoto E, Ueda Y, Kino T (1995) Purification of ultra-high purity aluminum. Suppl Phys II 5(C7):153–157 19. Lee HY, Oh JK, Dong HL (1990) Purification of tin by zone refining with development of a new model. Metal Trans B 21(3):455–461 20. Spimjr JA, Bernadou MJS, Garcia A (2000) Numerical modeling and optimization of zone refining. J Alloy Compd 298(1):299–305 21. Kurz W, Fisher D (1989) Fundamentals of solidification. Trans Tech Publications, Switzerland 22. Osono H, Maeta H, Matsusaka H (2005) Preparation of highly perfect aluminum crystal by cold-crucible induction melting in ultrahigh vacuum. Mater Trans 43(2):121–124 23. Kim P, Mihara Y (2000) high purity metal production using dry refining processes. Mater Trans JIM 41:37–43 24. Nakamura M, Watanabe M, Tanaka K (2014) Zone refining of aluminum and its simulation. Mater Trans 55:664–670 25. Cheung N, Bertazzoli R, Garcia A (2008) Experimental impurity segregation and numerical analysis based on variable solute distribution coefficients during multi-pass zone refining of aluminum. J Cryst Growth 310:1274–1280 26. Spim JA, Bernadou MJS, Garcia A (2000) Numerical modeling and optimization of zone refining. J Alloy Compd 298:299–305 27. Sen W, Fang HS, Jin ZL (2014) Integrated analysis and design optimization of germanium purification process using zone-refining technique. J Cryst Growth 408:42–48

Author Index

A Aadli, Ahmed S., 39 Abbasi, M., 579 Abdul Samad, Mohammed, 865 Abu Jadayil, Wisam, 413 Abu-Lebdeh, Taher, 355 Adegbuyi, Olasunkanmi B., 77 Agnew, Sean, 1383 Ahsan, Faiyaz, 319 Aigbodion, Aireguamen, 155 Ajayi, Oluseyi O., 667 Akinlabi, E.T., 279, 667 Akinlabi, S.T., 667 Aksu, Büsra, 1281 Alexander, S., 799 Alkan, Murat, 1281 Allison, Paul G., 1383 Alqarni, Laila, 1589 Alves, André L.M., 1459 Aminirastabi, Habibollah, 979 Amirkhiz, Babak Shalchi, 331, 443, 455 Anderson, I.E., 1507 An, Haifei, 1045 Antony, Veena, 819 An, Xinglong, 433 Arregui-Mena, J. David, 901, 907 Arthur, N., 279 Åsebø, Olav, 503 Asrar, Nausha, 927 Assis, Weslley L.S., 1459 Azar, Amin S., 503 Azizi, Hossein, 311 Azmah Hanim, M.A., 645

B Badgley, Peter, 519 Bagherpour, Iman, 143 Bai, Chenguang, 845 Baker, Paul A., 819 Bakhshinejad, Ali, 339 Bao, Yanping, 997 Baroutaji, A., 755 Barrado, Francys, 519 Bartoli, T., 607 Basu, Rahul, 879 Bayless, Trenin K., 1521 Beck, Jan, 887 Beckmann, Tobias, 1139 Behrens, Bernd-Arno, 719 Bein, Thilo, 371 Belt, Cynthia, 1079 Bergs, Thomas, 1615 Bessais, L., 593, 607 Bhattacharyya, Jishnu J., 1383 Bhatt, Dhananjay V., 697, 733 Bird, David, 1563, 1575, 1629 Biswas, Indrani, 1499 Bleck, Wolfgang, 1401 Bobzin, Kirsten, 1615 Bodude, Muideen Adebayo, 77 Boostani, Alireza Fadavi, 1383 Bordbari, Kambiz, 121 Bouzidi, W., 607 Brice, David, 599 Brögelmann, Tobias, 1615 Brøtan, Vegard, 503 Brown, A.D., 799 Burchell, Timothy D., 901

© The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6

1707

1708 C Cai Li, Ru, 1357 Cai, Zeyun, 1111 Cao, Xinjin, 205 Caravaca, Elbert, 1575, 1629 Carpenter, John, 197 Carton, J.G., 755 Casari, Daniele, 1665 Chandrasekar, Srinivasan, 599, 1499 Chavez, Luis, 887 Cheepu, Muralimohan, 259 Cheng, Hongwei, 1037 Cheng, Yi, 347 Chen, Xi, 549 Chen, Xudong, 1251 Chen, Zheng, 347 Che, Woo Seong, 259 Choi, Wan-Ho, 115 Chow, Tien See, 1589 Church, Benjamin, 339 Contescu, Cristian, 901 Couvrat, M., 1057 Cubides, Yenny, 1099 Cui, Hao, 137, 495 Cui, Xiaokang, 1111 Cutolo, Antonio, 395 D D’Aberle, Mark, 1521 Dai, Yongnian, 1697 Damptey, Ransford, 355 Daroowalla, Rayan, 1589 de Andrade, Lidiane Maria, 837 de Carvalho, Mariana Alves, 837 de Formanoir, Charlotte, 395 Dharmendra, Chalasani, 443 Dinda, Soumitra Kumar, 239 Dokumaci, Esra, 1281 Dong, Jian, 687 Donmez, M. Alkan, 269 dos Santos Mangualde, Cesar Bertolin, 637 Dou, Bingsheng, 687 Downey, Jerome P., 1521 Drouven, Carsten, 1401 Du, Jing, 793 Du, Shuang, 1589 Dvivedi, Akshay, 745 Dzisah, Patrick, 1545 E Easton, Mark, 1687 Edmondson, Philip D., 901 Eker, Emel, 1563

Author Index Elangeswaran, Chola, 395 El-Aziz, A.M., 185 El Kadiri, Haitham, 1383 Erwee, Markus Wouter, 1161 Espinosa, Denise Crocce Romano, 837, 967 Espinoza, Cesar, 1099 F Fang, Chao, 1483 Fan, Haidong, 1271, 1305 Fan, Zhou, 1483 Fellah, Mamoun, 865 Fen, Yunli, 1427 Fracchia, Elisa, 711 G Gallegos Pérez, A.I., 1389 Gao, Jingbao, 1027 Geng, Lan’xin, 347 Gerard, C., 1347 Ghamarian, Nima, 645 Gholipour, Javad, 205 Ghosh, Somnath, 657 Giorla, Alain B., 907 Giri, Anit, 1411 Gobber, Federico Simone, 711 Gong, Chujie, 793 Grau, Henry, 1629 Grundmeier, Guido, 485 Gunnarsson, C.A., 799 Guo, Chunhuan, 251 Guo, Haiding, 765 Guo, Hui, 1357 Guo, Wei, 1067 H Haase, Christian, 301 Hadadzadeh, Amir, 331, 443, 455 Hadi, Abdul-Sommed, 781 Ham, Jeffrey, 927 Hamoush, Sameer, 355 Hansen, Samuel, 289 Han, Xu, 63 Han, Yun, 571 He, Feiyu, 687 Heigel, Jarred C., 269 Henderson, Kevin, 197 Hernandez, F., 1507 Hezil, Naouel, 865 Hild, Rafael, 1615 Hill, Bryce E., 781 Hitzler, Leonhard, 407 Hofer, Dominik, 1139

Author Index Hoffmann, Dennis C., 1615 Hogg, Ben, 1149 Honda, Yoshiaki, 529 Hong, Tae-Hyun, 115 Horan, Caleb, 197 Horie, Tetsuro, 827 Hossain, Samiha, 1643 Hossaini-Zadeh, Mehran, 793 Hosseini Far, A.R., 579 Hou, Da Wei, 1067 Hoyer, Kay-Peter, 475, 485 Huang, Bensheng, 1483 Huang, Hongtao, 687 Hübner, Sven, 719 Hu, Songhao, 1263 I Ifijen, Hilary I., 155 Ikhuoria, Esther U., 155 Ilyas, Muhammad Tasaduq, 681 Imam, Muhammad A., 1471 Inegbenebor, A.O., 667 Ingole, Sudeep P., 697, 733 Iost, Alain, 865 J Jackman, Jennifer, 1687 Jellison, G.E., 907 Jiang, Fengchun, 251 Jiang, Guangrui, 949 Jiang, Wenhuan, 63 Ji, Gouli, 979 Joseph, Olufunmilola, 959 Joubert, Hugo, 1181 Jun, Kai-Feng, 165 K Kainer, K.U., 1677 Kara, Aslihan, 1281 Karimidehcheshmeh, Fatemeh, 979 Katoh, Yutai, 901 Kawai, Masaki, 907 Kawata, Hiroyuki, 529 Keßler, Olaf, 485 Kilymis, D., 1347 Knapp, Cameron, 197 Koike, Mari, 827 Kong, B., 1507 Kong, Lingxin, 1027 Kossman, Stephania, 865 Kreitman, Meir, 269 Krivanec, Cory, 1383 Kruppe, Nathan C., 1615

1709 Kumar, Nilesh, 1089 Kumar, Pradeep, 745 Kundin, Julia, 301 Kurz, G., 1677 L Ladani, Leila, 319 Lamberti, Vincent, 355 Langelier, Brian, 331 Lan, Peng, 51, 93 Laquidara, Joseph, 1575 Lee, Jung-Hoon, 115 Le Pape, Yann, 907 Leramo, Richard, 939 Letzig, D., 1677 Ley, Nathan A., 1411 Liang, Jinglong, 1199, 1227 Liao, Yilong, 845 Libera, Ondřej, 1325 Li, Changrong, 549 Li, Dianzhong, 3 Liewald, Mathias, 1603 Li, Guangshi, 1037 Li, Hong-Yi, 165 Li, Hui, 1199, 1227 Li, Hui-cheng, 1437 Li, Jian, 331, 455 Li, Jie, 17, 1045, 1357 Li, Jingshe, 29, 421, 561 Li, Jun-Guo, 1217, 1237 Li, Liang, 51, 93 Li, Longfei, 1111 Lim, H.N., 645 Lin, Yirong, 887 Li, Qingjuan, 855 Li, Shenggang, 1037 Lisiecka-Graca, Paulina, 537 Li, Tianshuang, 1357 Liu, Baicheng, 1007 Liu, Fangzhou, 1263 Liu, Fei, 433 Liu, Guanghui, 949 Liu, Huasai, 571 Liu, Kun, 1427 Liu, Meihuan, 63 Liu, Roulei, 1589 Liu, Shaojun, 433 Liu, Shiyuan, 495, 845, 855 Liu, Shuai, 549 Liu, Shun Ming, 1067 Liu, Wei, 29 Liu, Weixing, 17, 1045 Liu, Xiaogang, 765

1710 Liu, Yan, 1589 Liu, Yang, 1447 Liu, Yingli, 173 Liu, Yong, 1263 Liu, Yu-xiang, 1437 Liu, Zhaoyang, 463 Liu, Zhen, 1111, 1437 Liu, Ziyi, 997 Li, Wei, 1271 Li, Xiangyu, 571 Li, Xiao-Yu, 1217 Li, Yang, 1335 Li, Yang Long, 1067 Li, Yanjun, 1665 Li, Yifu, 1027 Li, Yungang, 915, 1199 Li, Yusheng, 1251 Li, Zhiye, 657 López Martínez, E., 1389 López Soria, J.J., 1389 Loto, Cleophas Akintoye, 959 Loto, Roland Tolulope, 939, 959 Lu, Feng, 63 Luhmann, Keith, 1575 Luidold, Stefan, 1139 Luo, Xia, 1483 Lu, Tengfei, 495, 855 Lu, Xionggang, 1037 M Maaza, Malik, 155 Magne, D., 1057 Mahajanam, Sudhakar, 1099 Majta, Janusz, 537 Mannens, Robby, 1615 Mann, James B., 599, 1499 Martin-Root, Chris, 519 Masina, B., 279 Mathiesen, Ragnvald H., 1665 Mathoho, I., 279 Mattfeld, Patrick, 1615 Ma, Xianfeng, 1335 Mc Dougall, Isobel, 1181 Mejias, Alberto, 865 Melz, Tobias, 383 Menghani, Jyoti V., 697, 733 Merkel, Markus, 407 Michalowicz, A., 607 Mirkouei, Amin, 289 Mitchell, Richard J., 827 Mliki, N., 607 Mohammadi, Mohsen, 311, 331, 443, 455 Möhwald, Kai, 719

Author Index Möller, Benjamin, 383 Montagne, Alex, 865 More, Satish R., 697 Moscovici, J., 607 Mousavi Anijdan, S.H., 579 Mujahid, Shiraz, 1383 Muralidharan, Gokula Krishna, 395 Murr, Lawrence, 887 Murty, Korukonda L., 1089 Muszka, Krzysztof, 537 Muthupandi, V., 259 Mypati, Omkar, 619 N Nadeem, Amer, 681 Nahavandi, M., 645 Nakouri, K., 593 Nemir, David, 887 Nezafati, Marjan, 339 Ni, Guolong, 1427 Nikolic, Stanko, 1149 Ni, Song, 433 Nnaji, Ruth Nkiruka, 77 O Obeidi, M., 755 O’Brien, Daniel J., 657 Öchsner, Andreas, 407 Oghenevwegba, Emuowhochere, 667 Ohba, M., 967 Okabe, Toru, 827 Olabi, A.G., 755 Oladoye, A.M., 755 Olaitan, Akanji, 959 Omorogbe, Stanley O., 155 Oppedal, Andrew L., 1383 Otto, Robert, 503 Ourdjini, Ali, 645 Overby, David, 519 Oyebade, Babatunde, 939 Özkaya, Fahrettin, 719 P Pakulat, S., 1677 Pal, Surjya Kanta, 619 Pal, Tapan Kumar, 217 Panchal, Viral, 1629 Panchhi, Navjot, 1589 Panossian, Z., 967 Park, Jin-Seong, 115 Patterson, Brian, 197 Paul-Boncour, V., 593 Peng, Yingbo, 1263

Author Index Peng, Yuxiang, 1313 Peng, Zheng-Wu, 165 Pizzagalli, L., 1347 Plaisted, T.A., 799 Polák, Jaroslav, 1125 Prahl, Ulrich, 301 Proano, Camila, 1207 Provatas, Nikolas, 311 Puentes Rodriguez, Brhayan Stiven, 599 Q Qiu, Guibao, 495, 845, 855 Qiu, Wanling, 949 R Rafaels, K.A., 799 Ramazani, Ali, 301 Ramver, 745 Rathod, M.J., 733 Ravindra, Nuggehalli M., 1545, 1563, 1575, 1589, 1629, 1643 Raza, Mohsin Ali, 681 Reddy, Ramana G., 809, 1227, 1313, 1471 Regele, J.D., 1507 Reichardt, Gerd, 1603 Reyes, Rodrigo Valenzuela, 637 Reynolds, Quinn Gareth, 1161 Riedemann, T., 1507 Rios, Paulo R., 1459 Robertson, Christian, 1335 Rohatgi Pradeep, 339 Rosario, Carlos Gonzalo Alvarez, 837 Rosseel, Thomas M., 907 Rosso, Mario, 711 Roy, Gour Gopal, 239 S Sadayappan, Kumar, 1687 Sadoh, Airefetalo, 1545 Salawu, Enesi Y., 667 Salwén, Anders, 1019 Sanders, Paul, 1291, 1421 Sarkar, Tapan, 217 Sauvage, X., 1057 Scarazzato, T., 967 Schaper, Mirko, 475, 485 Schneider, Judith, 197 Schöler, Simon, 719 Schulenburg, Frank, 1139 Scurria, Matilde, 371, 383 Serhan, Duaa, 413 Sert, Enes, 407

1711 Shang, Ting, 949 Sheng, Jiazhen, 115 Shi, Shengyun, 63 Shivaram, Vishwas Danthi, 1589 Siahooei, Mohsen Ameri, 121 Simo, Aline, 155 Smith, Jesse, 1411 Song, Bo, 1111 Song, Chunyan, 1427 Song, Lijun, 463 Song, Min, 433 Song, Wenwen, 1401 Spinelli, José Eduardo, 637 Srba, Ondřej, 1325 Srirangam, Prakash, 239, 619 Steenkamp, Joalet Dalene, 1161 Stockman, Tom, 197 Stokes, J., 755 Sun, Laibo, 251 Sun, Mingyue, 3 Sun, Xiaojing, 251 Sun, Yan-hui, 687, 1447 Surolia, Ranu, 819 Su, Yan, 251 Švrčula, Petr, 1325 Swanepoel, Stefan, 1161 T Tajuelo-Rodriguez, Elena, 907 Takeda, Kengo, 529 Tanaka, Takaaki, 513 Tang, Haiyan, 51, 93 Tang, Jing, 1271, 1305 Tarek, M., 185 Tasche, Lennart, 485 Tavangarian, Fariborz, 1207, 1533 Tenório, Jorge Alberto Soares, 837, 967 Thomas, Vinoy, 819 Tian, Can, 1045 Tiarks, J., 1507 Tie, Zhanpeng, 51, 93 Tjayadi, Leonardi, 1089 Tlotleng, M., 279 Toji, Yuki, 513 Torrence, Christa E., 907 Touhami, Mohamed Zine, 865 Trauth, Daniel, 1615 Trumble, Kevin, 599, 1499 Tsuchida, Noriyuki, 513 Tucker, Bernabe S., 819 Türkoglu, Berkay, 1281 Twomey, B., 755

1712 U Ugurluer, Dilan, 1281 V Van Hooreweder, Brecht, 395 Vázquez Gómez, O., 1389 Ventura, Harison S., 1459 Vergara Hernández, H.J., 1389 Villa, Elena, 1459 Vohra, Yogesh, 819 Voigt, Paul, 1149 W Wagener, Rainer, 371, 383 Wagh, S.V., 733 Wallace, Grant C., 1521 Wang, Biao, 1335 Wang, Bo, 433 Wang, Cong, 915 Wang, Dongbin, 1199, 1227 Wang, Feng, 29 Wang, Fuming, 549 Wang, Hui, 1067 Wang, Jian, 495, 845 Wang, Jing, 1227 Wang, Jingkun, 1357 Wang, Jingsong, 173 Wang, Min, 997 Wang, Qingyuan, 1271 Wang, Qiyu, 1357 Wang, Rongyue, 105 Wang, Shuhuan, 1427 Wang, Shu-sen, 629, 1373 Wang, Yajie, 173 Wang, Ya-Jun, 1217, 1237 Wang, Yueping, 205 Wan, Heli, 1697 Wanjara, Priti, 205 Ward, T., 1507 Weaver, Jordan S., 269 Weerasooriya, T., 799 Wei, DeAn, 1305 Wen, Liangying, 63 Werner, Ewald, 407 Wheeler, Robert W., 1411 Whittington, Wilburn R., 1383 Worth, Robert N., 901 X Xie, Bing, 165 Xie, Chunqian, 571 Xiong, Xiaolu, 1037

Author Index Xi, Xiaojun, 421, 561 Xu, Baoqiang, 1027, 1697 Xu, Bin, 3 Xu, Junjie, 1027 Xu, Qian, 1037 Xu, Qingyan, 1007 Xu, Yijiang, 1665 Y Yang, Aimin, 17, 1045 Yang, Bin, 1027, 1251, 1697 Yang, Cong, 1007 Yang, Jun, 1483 Yang, Shufeng, 29, 421, 561 Yang, Yang, 1421 Yang, Yu, 1227 Yao, Da-wei, 629, 1373 Yao, Yu, 347 Ye, Maolin, 421, 561 Yilkiran, Deniz, 719 Young, Jacob, 809 Young, Marcus L., 1411 Yuan, Ding, 251 Yuan, Zhangfu, 105 Yu, Meng, 1067 Yu, Xiangtao, 105 Z Zainal, Zulkarnain, 645 Zaunschirm, Stefan, 1325 Zeng, Ya-Nan, 1217, 1237 Zhai, Qi-jie, 1437 Zháňal, Pavel, 1325 Zhang, GuoZhi, 137 Zhang, Hongliang, 1357 Zhang, Honglin, 3 Zhang, Jiaquan, 51, 93 Zhang, Jieyu, 347 Zhang, Jun, 137 Zhang, Kaixuan, 1427 Zhang, Mingyang, 1263 Zhang, Rich, 519 Zhang, Sarah, 519 Zhang, Shixian, 915 Zhang, Wanlong, 173 Zhang, Wei, 1263 Zhang, Xiaofan, 657 Zhang, Xiaoxing, 137 Zhang, Xu, 1305 Zhang, Xueliang, 29 Zhang, Yuan-wang, 629, 1373 Zhang, Yuzhu, 17

Author Index Zhao, Jinyang, 1697 Zhao, Xiao, 51, 93 Zhao, Xiaoping, 915 Zhen, Jianping, 687 Zhou, Tihe, 519 Zhou, Yuxiao, 793 Zhu, Libin, 29, 561 Zhu, Qiang, 463

1713 Zhuravlev, Evgeny, 485 Zhu, Yuntian, 1251 Zimina, Mariia, 1325 Zolko, Caleb, 1207, 1533 Zou, Bowen, 1401 Zuo, Haibin, 173 Zuo, Yongji, 765

Subject Index

Numbers 316L, 143, 145, 149–151, 153, 154, 219, 395–400, 402–404, 683, 688, 699 316L steel, 143, 145–148, 151, 153, 280 3D printers and its role in printing gears, 1593 3D printing and off-the-shelf interface, 292 3D printing for pharmaceutics and transdermal drug delivery, 1565 3D printing, 292, 1207–1209, 1563, 1565, 1567, 1568, 1570, 1572, 1575, 1589, 1593, 1594 A Ab initio molecular dynamics, 3, 4 Activation energy, 620, 866, 998, 1045, 1046, 1051, 1053, 1155, 1389, 1390, 1397, 1398 Activation energy of austenite, 1397 Activity, 107, 110, 144, 422, 831, 849, 927, 1028–1032, 1034, 1114, 1130, 1231, 1293, 1296, 1327, 1340, 1348, 1443, 1473, 1476–1478, 1567, 1688 Additive manufacturing, 197, 206, 251, 252, 269, 270, 279, 289, 319, 320, 331, 332, 339, 355, 372, 374, 383, 392, 395, 434, 443, 444, 449, 452, 455, 456, 476, 486, 492, 503, 504, 804, 827, 828, 835, 1207–1209, 1214, 1564, 1575, 1594 Adsorption gas, 122 Advanced high-strength steels, 519, 529, 572, 1390 Advantages of Contactless Gears, 1593

Akermanite, 1533–1537, 1539, 1540 AlCoCrCu0.5FeNi, 1271–1278 Alloy and simulation conditions, 1010 AlN precipitation, 999, 1001, 1003, 1004 AlSi10Mg, 331–337, 339, 340, 342–345, 383, 384, 392, 408, 409, 490, 491 Aluminium, 4, 39, 41, 42, 44, 45, 121, 240, 243, 245, 407, 408, 622, 755, 756, 1200, 1358, 1401, 1625, 1665, 1688 Aluminium alloy, 240, 699, 1502 Aluminum, 40, 41, 44–48, 121, 122, 126, 130, 205, 215, 260, 336, 339, 384, 422, 448, 485, 486, 488, 490, 600, 619, 620, 647, 734, 755, 756, 765, 766, 772, 801, 822, 828, 840, 843, 950, 997, 998, 1079, 1081, 1119, 1199–1201, 1204, 1229, 1232–1234, 1272, 1282–1284, 1294, 1298, 1299, 1357, 1358, 1362–1366, 1368, 1422, 1423, 1439, 1595, 1665, 1666, 1688, 1697, 1698, 1701, 1704 Aluminum alloys, 332, 485, 486, 490, 492, 756, 1127, 1421, 1693 Amorphous metal, 887 Amorphous silicon, 1347, 1413 Analysis of MgO’s influence on pellet performance, 22 Analysis of surface film, 6 Analysis of temperature variation in molten pool, 108 Analysis of the influence of shot peening parameters on evolving surface textures, 1619

© The Minerals, Metals & Materials Society 2019 The Minerals, Metals & Materials Society (ed.), TMS 2019 148th Annual Meeting & Exhibition Supplemental Proceedings, The Minerals, Metals & Materials Series, https://doi.org/10.1007/978-3-030-05861-6

1715

1716 Analysis of the sample section, 1201 Analysis of Ti6–Al4–V aging phase transition process, 773 Analytical determination of the averaged, projected surface normal pressure, 1610 Anisotropy, 271, 274, 332, 371, 384, 388, 408, 410, 421, 422, 427–429, 1442 Annealing, 410, 456, 519, 520, 522–526, 571, 573–576, 600, 603–605, 629, 630, 632, 788, 960, 965, 1067, 1068, 1264, 1272, 1374, 1378, 1404–1407, 1429, 1533, 1535–1540 AOD slag, 1217, 1218, 1220 Applications, 48, 78, 122, 138, 143, 144, 155, 156, 160, 162, 185, 186, 197–199, 206, 218, 239, 259, 260, 270, 290, 339, 365, 366, 396, 413, 434, 443, 455, 456, 482, 495, 496, 504, 538, 579, 580, 600, 637, 638, 645, 647, 681, 682, 698, 700, 712, 734, 746, 756, 782, 783, 787, 809, 816, 819, 820, 845, 846, 866, 888, 889, 908, 911, 1037, 1057, 1100, 1101, 1140, 1152, 1157, 1172, 1184, 1263, 1269, 1281, 1314, 1325, 1330, 1347, 1348, 1411, 1412, 1416, 1419, 1484, 1521, 1567, 1568, 1571, 1576, 1578, 1589, 1590, 1593, 1594, 1599, 1643, 1687–1689, 1691, 1692, 1695 Area fraction transformed against time, 1463 Artificial intelligent and grain growth, 991 As-cast structure, 602 As-welded microstructure, 225 Atomic oxygen adsorption on the (001) surface, 1039 Atom Probe Tomography (APT), 331–333, 336, 337, 1402, 1404, 1407 Au–Cu alloy, 1027, 1028, 1031–1034 Austempered ductile iron, 217, 218 Austempered microstructure, 226 Austenite formation, 1392 Austenite formation rate, 1395 Austenitic stainless steel, 107, 218, 219, 929, 1089, 1090, 1129, 1251 Autonomous self-healing, 1645, 1646 B Babble fluidized bed, 63, 76 Bacterial adherence and Biofilm formation on alloy specimens, 830 Bacterial suspensions, 830 Ball-milling technology, 496, 845 Base metal microstructure, 224 Basicity, 1046, 1218

Subject Index Basic principles of molecular dynamics, 767 Batch leaching test, 1219 Battery, 171, 619, 620, 838, 1348, 1594–1596 Bead geometry, 323, 326 Beam oscillation, 239–241, 245–247 Bending test, 1265 Bi-Ag, 646–648, 652 Bicrystals, 1306, 1309 Biodegradable, 809, 810, 816, 817 Bioenergy production, 289 Biofilm, 827, 828, 830–832, 834, 835 Biofuel, 290, 296 Bioleaching, 837–843 Bioluminescence assay, 827, 832 Biomass, 289, 290, 293, 296 Biomaterials, 810, 819, 820, 856, 1484, 1570, 1572, 1575, 1576 Biomedical applications, 475, 476, 810, 846 Bone, 1646 Bone surrogates, 801, 803 Boundary condition and physical property, 33 Boundary conditions, 52, 73, 767, 768, 771, 1019–1021, 1071, 1168, 1272, 1306, 1461, 1603, 1617 Boundary layer, 720, 879, 881, 883, 885, 1048, 1049, 1051, 1704 Brittle temperature range, 96 Bulk and surface properties of NiS, 1039 Bulk Metallic Glass (BMG), 879, 880, 888, 893, 897 Butler–Volmer equation, 967, 969, 971, 974 C CaF2 electrolyte, 1475 Calculation of decarburization reaction, 110 Calculation of geometrical necessary dislocation density based on the EBSD data, 541 Capstone, 1291, 1292, 1299, 1300 Carbide-free bainite, 550 Carbides, 227, 425, 525, 553, 575, 712–715, 717, 725, 960, 1057–1065, 1101, 1103, 1116, 1119, 1405 Carbon nanotube, 121, 122, 126–130, 133 Catalytic Fast Pyrolysis (CFP), 290 CATME Peer Evaluation, 1298 Cemented carbide, 711–713, 715 Cementitious material, 1217, 1219–1224 Ceramics, 1658 Cetyltrimethylammonium bromide, 155, 157, 162 Characterization, 4, 8, 13, 143, 146, 151, 157, 166, 175, 186, 189, 193, 198, 239–241,

Subject Index 259, 261, 280, 281, 331, 399, 487, 689, 724, 745, 747, 783, 812, 839, 865, 867, 868, 907, 909, 917, 940, 960, 1089, 1091, 1093, 1094, 1330, 1413, 1474, 1579, 1603, 1612, 1622, 1630, 1666 Characterization measurements, 917 Characterization techniques, 157 Charge analysis, 3, 12, 1037, 1041 Charpy impact, 231 Charpy impact test, 221, 224, 232, 583 Chemical analysis, 582 Chemical composition of weld deposits, 219 Chemicals, 157 Chemistry design, 520 Chromium leaching, 1218, 1220–1223 Classifying of composites, 79 Coal bottom ash, 697, 698, 700, 702–704, 708, 709 Coating, 143–148, 151–155, 174, 218, 219, 261, 280, 600, 681–684, 687–695, 720, 721, 755–758, 760, 761, 783, 814, 916, 919–925, 927–937, 949–956, 967, 1199–1205, 1610, 1615, 1618, 1621, 1624, 1626, 1633, 1644, 1656, 1657 Coating agent, 1656 Cobalt deposition, 967–969, 971, 972, 975 CoBlast, 756–758, 760, 761 Co–Cr–Mo–W alloys, 433, 435–440 Coefficient of friction, 728 Coils, 522, 1598, 1599 Cold forming, 572, 1615, 1616 Cold rolled, 571, 572, 599, 601, 950, 1390, 1413–1417, 1419 Cold rolling, 521, 522, 599–601, 603–605, 1067, 1068 Collaborations, 1692 Colour reflectance, 161 Comminution, 839 Compaction pressure, 856, 857, 861, 863 Comparison of flat and deflected strip drawing experiments/discussion of deviations, 1610 Comparison of grain growth between simulations, 988 Comparison of the results, 1611 Composite, 78–89, 91, 144, 173, 174, 176, 177, 179–182, 208, 339, 344, 345, 495, 496, 498, 500, 689, 694, 695, 712, 800, 804, 807, 911, 915, 917, 919–921, 924, 1217, 1219–1224, 1363, 1648, 1653, 1655, 1656

1717 Composition of iron ore powder phase, 18 Composition, melting point, specific heat capacity, and enthalpy of fusion of freeze-lining, 1170 Compressive properties, 80, 173, 174, 177 Compressive properties of composites, 177 Compressive stress at break against untreated, 85 Compressive stress at yield against untreated samples, 86 Computational materials, 1292, 1294, 1299, 1300, 1691 Computer-Aided Design (CAD), 292 Computer simulation, 356, 980, 1460, 1461 Computer simulation methodology, 1461 Concrete, 79, 365, 366, 907–911, 1090, 1218, 1658 Consolidation results, 892 Contact area analysis, 795 Contactless gears, 1593, 1599 Contiguity, 1465 Continuous annealing furnace, 1067, 1074 Continuous casting, 29, 51–53, 60, 93, 94, 96, 102, 563 Continuous heating, 1389, 1390, 1392, 1393 Continuous improvement, 1079, 1085 Control equation, 31 Convective conditions, 881 Converting, 888, 897, 1150, 1151, 1157, 1182, 1193 Convex roll soft reduction, 51, 93, 94, 102 Copper, 4, 10, 187, 189, 212, 215, 260, 350, 413–415, 417, 418, 452, 601, 619, 620, 622, 629, 630, 643, 647, 681, 683, 699, 747, 749, 766, 810, 814, 837–843, 881, 889, 935, 1028, 1125–1129, 1132, 1150, 1151, 1153, 1154, 1156, 1157, 1181, 1185, 1186, 1192, 1282, 1283, 1307, 1309, 1330, 1373, 1374, 1376, 1380, 1424, 1666, 1701 Copper cooling, 1149, 1158, 1182, 1185, 1186, 1191–1194 Copper recovery, 838 Correction function, 1466 Corrosion, 143, 145, 147, 151–154, 174, 206, 213, 332, 448, 455, 475, 476, 478, 481, 492, 561, 562, 564–568, 681–684, 698, 700, 755–761, 809–811, 814–817, 838, 840, 887, 915, 916, 920–923, 927–929, 932–934, 936, 939–944, 947, 949–954, 956, 959–965, 1058, 1068, 1089–1091,

1718 1093–1095, 1100, 1101, 1182, 1199–1205, 1251, 1252, 1255, 1257, 1258, 1294, 1298, 1313–1320, 1326, 1329, 1331, 1332, 1427, 1447, 1484, 1490, 1604, 1656, 1658, 1687, 1693 Corrosion analysis, 478 Corrosion behavior, 143, 145, 561, 562, 638, 681, 810, 949, 1484 Corrosion resistance, 143, 144, 152, 154, 173, 174, 259, 280, 414, 421, 434, 443, 456, 464, 487, 561–563, 565–568, 638, 681–684, 758, 760, 814, 846, 855, 888, 915–917, 922, 923, 925, 927, 928, 933, 949, 950, 959–961, 965, 1028, 1057, 1100, 1101, 1107, 1112, 1199–1205, 1251, 1252, 1255, 1258, 1259, 1263, 1271, 1281, 1326, 1448, 1484, 1698 Cost-effectiveness and scalability, 1571 Crack growth, 1089–1096, 1103, 1104, 1125–1127, 1130–1134, 1327, 1328 Crack growth during stress corrosion cracking, 1093 Crack initiation, 232, 379, 380, 387–390, 392, 393, 395, 396, 398, 401, 402, 404, 409, 410, 714, 1092, 1126, 1127, 1129, 1131–1133 Cranial bone, 799–801, 804, 806, 807 Cr6+ concentration, 1219–1221, 1223 Creep, 53, 79, 464, 503, 888, 902, 906, 1057, 1058, 1062, 1063, 1065, 1100, 1101, 1133, 1326, 1384, 1421, 1490, 1648, 1657, 1693 Critical size for bubble passing interface, 37 Cr–Mo steel, 1100, 1107, 1108 Crystal selector, 348 Current Course Content, 1293 Cu–0.3wt%Ag, 629, 630, 632–634 Cycles to crack initiation, 389 Cyclic material behavior, 371, 372, 374, 377, 384, 388, 389, 392 Cyclic stress–strain behavior, 373, 377, 378, 383, 384, 387, 392 Cyclosizer particle diameter separation, 1523 D Data analysis issues, 1084 Decarbonization and chromium conservation, 105 Deep drawing, 730, 1603, 1604 Defect-Induced Apparent Temperature (DIAT) shift, 1335–1343

Subject Index Deformation model, 53 Deformation twinning, 1271, 1272, 1274–1277 Degradation, 19, 291, 475, 476, 478, 479, 481, 482, 698, 810, 820, 902, 908, 939, 960, 961, 1090, 1112, 1119, 1252, 1325, 1533, 1643, 1648, 1649 Degree of polymerization, 1357, 1366 d-ferrite, 1447–1451, 1456 d-ferrite phase melting and solidification in DSS, 1450 d!c Phase Transformation in DSS, 1449 Dendrite growth, 1007, 1008, 1010–1012, 1014–1016 Dendritic morphology, 1437, 1438, 1440, 1441, 1443, 1669 Density Functionals Theory (DFT), 4, 1042, 1359 Density separation, 1522 Density, specific heat capacity, and thermal conductivity of steel shell and refractory materials, 1169 Dental alloys, 827, 828 Dental implant, 794 Deposition process of the (Cr,Al)N + Mo:S coating, 1618 Details of molecular dynamics simulation, 1359 Detection of inclusions, 563 Determination of kd, 364 Development and use of a pin-on-cylinder tribometer, 1620 Development of heat transfer model, 1167 DFSC has further been extended by the Differential Reheating Method (DRM), 490 Diamond-Like Carbon (DLC) coating application, 929 Differential Fast Scanning Calorimetry (DFSC), 489 Differential Scanning Calorimetry (DSC), 158, 160, 177, 178, 822–824, 1143, 1268, 1411, 1412, 1414, 1486, 1491 Different nucleation mechanisms, 1002 Diffusion bonding, 4, 260, 261, 765–768, 774–776 Diffusion coefficient, 426, 620, 765, 771, 776, 879, 979, 1009, 1046, 1111, 1112, 1117–1119, 1121, 1241, 1360, 1361, 1363, 1366, 1368, 1491, 1492, 1700 Diffusion-controlled, 1019, 1021, 1022 Digital twin, 371–373, 376–378

Subject Index Digital twin for fatigue approach, 372 Digital volume correlation, 793–795 Dilatometric Analysis, 1391 Dilatometry, 1391, 1394 Direct current potential drop technique for crack size measurement, 1091 Directional solidification, 1670 Direct Metal Laser Sintering (DMLS), 280, 331–337, 455–461 Dislocation density, 255, 269, 273, 276, 277, 530, 538, 541–543, 545, 546, 553, 589, 870, 1111, 1113, 1120, 1130, 1340, 1341, 1378 Dislocation dynamics, 1305, 1306, 1309, 1335, 1336, 1342 Dispersion of graphene, 175 Dissimilar Metal Weld (DMW), 1099–1108 Dissolution, 3, 4, 6, 8–10, 12–14, 41, 156, 336, 490, 521, 562, 566, 746, 747, 750–752, 921, 922, 1019, 1024, 1025, 1045–1053, 1155, 1170, 1175, 1177, 1217, 1221, 1223, 1224, 1231, 1232, 1258, 1267, 1268, 1390, 1404, 1490 Dissolution behavior, 3, 4, 1046 Distribution regularity of types of aluminum fluoride complex ions and fluorine ions, 1365 Drill collar, 928–934, 936 Dry metal forming, 720, 1617–1619 Ductile fracture, 232, 428, 429, 451, 452, 554, 648, 652 Ductile to Brittle Transition Temperature (DBTT) shift, 1336, 1338, 1339, 1343 Duplex Steel, 959, 960, 965 Dynamic Mechanical Analysis (DMA), 1582 E E36 shipbuilding steel, 421, 427 EBSD method, 546 Effect of ball-milling on raw materials, 848 Effect of carbon content on the reaction equilibrium, 109 Effect of compaction pressure on pore characteristics of porous titanium, 857 Effect of FeV80 on microstructure, 849 Effect of FeV80 on the porosity and mechanical property, 852 Effect of grain size, 532 Effect of laser power on hardness, 740 Effect of laser scanning speed hardness, 740

1719 Effect of MgO content on high-temperature properties of pellet, 24 Effect of MgO content on pellet bonding, 23 Effect of original Sm content on phase composition and morphology, 1431 Effect of PH and Temperature on carbonation, 42 Effect of pressing pressure on mechanical properties of porous titanium, 861 Effect of raw material size, 1232 Effect of surface textures on the punch force in an extrusion process, 1621 Effect of temperature on dissolution rate of silica particles, 1050 Effect of the power for ultrasonic micro-forging, 255 Effect of the temperature for ultrasonic micro-forging, 254 Effect of the time for ultrasonic micro-forging, 253 Effect of Y-based RE on the morphology, size and type of inclusions, 564, 565 Effect of Y on the microstructure, 425 Effects of a stress relief heat treatment, 387 Effects of solidification rate on component segregation, 1428 EH36 steel, 561, 564, 565, 567 Electrical conductivity, 165, 173, 174, 180, 181, 619, 620, 623, 624, 626, 630, 631, 633, 634, 682, 980, 1140, 1373, 1374, 1376–1378, 1380 Electrical conductivity study, 623 Electrical heating properties and heat transfer of composites, the, 180 Electrical steels, 600 Electric vehicles, 619, 1592, 1594–1596, 1599 Electric vehicles and their battery manufacturers, 1594 Electrochemical measurements, 167 Electrochemical results, 953 Electrodeposition, 261, 687–689, 693, 694, 838, 915–918 Electrode preparation, 1473 Electro healing, 1656 Electron Backscatter Diffraction (EBSD), 240, 270–275, 277, 331–335, 337, 348, 435, 455–461, 510, 531, 537, 538, 540, 541, 546, 551, 556, 557, 907, 911, 1266, 1267, 1383, 1385 Electron Beam Welding (EBW), 239–241

1720 Electron microscopy, 145, 147, 167, 174, 185, 189, 222, 239, 241, 435, 443, 444, 452, 456, 478, 490, 503, 531, 550, 551, 563, 573, 630, 726, 745, 747, 783, 811, 868, 887, 889, 903, 951, 1059, 1068, 1091, 1101, 1113, 1129, 1237, 1238, 1325, 1374, 1411, 1412, 1423, 1533, 1534 Electrophoretic, 143–145, 147 Electrophoretic deposition, 153, 681, 682, 756, 758 Electrospinning, 819–822 Electrothermal properties, 173 Elemental analysis, 811 Emissivity relative to energetic materials, 1630 Energetic materials, 1575, 1576, 1579, 1630 Energy, 5, 31, 78, 105, 107, 123–125, 141, 159, 165, 166, 170–172, 174, 206, 232, 239, 246, 252, 254, 255, 261, 262, 279, 281, 290, 291, 296, 320, 322, 333, 339, 340, 342, 343, 422, 423, 425, 435, 437, 439, 440, 448, 464, 467, 476–478, 486, 489, 491, 495–497, 504, 507, 522, 525, 538, 549, 550, 554, 557–559, 563, 573, 583, 584, 588, 623, 639, 681, 687, 689, 699, 722, 734, 735, 756–758, 767–769, 783, 845, 847–849, 853, 855, 857, 866, 869, 871, 897, 909, 921, 922, 935, 943, 946, 979–982, 986, 998–1003, 1008–1011, 1022–1024, 1029, 1037–1041, 1051, 1079, 1080, 1082–1086, 1091, 1093, 1102, 1114, 1118, 1140, 1151, 1158, 1162, 1167, 1172, 1181, 1183, 1187, 1231, 1233, 1237–1241, 1255, 1264, 1271–1274, 1276–1278, 1281, 1283, 1305, 1313, 1314, 1318–1320, 1325–1327, 1340, 1349, 1357, 1359, 1375, 1384, 1398, 1402–1404, 1411–1413, 1416, 1417, 1427, 1434, 1436, 1442, 1443, 1447, 1450, 1454, 1456, 1471, 1473, 1476–1478, 1502, 1505, 1525, 1530, 1569, 1576, 1593, 1595, 1597, 1599, 1629, 1630, 1648, 1649, 1652, 1688, 1691 Energy savings, 1079, 1080, 1146, 1147 Energy spectrum analysis, 1051 Environmental Considerations, Health and Safety, 1577 Equilibrium solidification, 1427, 1428, 1430, 1431, 1436, 1491

Subject Index Erosion, 111, 697, 698, 700, 703–706, 708, 709, 746, 927–932, 934–936, 1067, 1068, 1149 Erosion wear, 697–701, 704–706, 708, 709 Erosion wear behaviour, 705 Essential oil, 947 Establishment of dynamic model, 1048 Establishment of kinetic parameters, 1051 Establishment of weld procedure, 221 Evaluation of biological activity for alloy specimen, 831 Evaluation of experimental results, 1607 Evaluation of friction coefficient, 871 Evaluation of volume and wear rate, 872 EV battery manufacturers, 1594 EV manufacturers market share, 1595 Evolution of size and quantity of precipitates in solute-rich and solute-depleted regions, 1245 Evolution of the SUSEN hot cells, 1326 Excavation, 1161, 1162, 1178 Experimental parameters, 703 Experimental setup of the deflected strip drawing testing rig, 1605 Experimental study of feeding speed, 70 Extreme operating conditions and furnace cooling design, 1191 Extrusion-based bioprinting and fused deposition modelling, 1566 F Fatigue, 217–219, 221, 224, 232–235, 280, 282, 347, 371–378, 383–386, 388, 389, 391, 392, 395–398, 400–404, 407–410, 413–418, 421, 464, 492, 562, 700, 711, 714, 716, 927, 928, 1091, 1092, 1094–1096, 1100, 1125, 1126, 1128, 1131–1134, 1271, 1326, 1384, 1593, 1648, 1657, 1658 Fatigue life, 232, 233, 372, 373, 376, 377, 380, 389, 392, 409, 413–418, 1125, 1132–1134 Fatigue life testing, 415 Fatigue pre-cracking, 1091–1093 Fatigue properties, 400 Fe–Al alloy, 687–691, 694, 695 Feasibility analysis, 107 FeCrNi alloys, 1057–1064 Feeding speed, 69 Ferrovanadium, 845

Subject Index Fibroblasts, 821, 822 FIB-SEM tomography, 901, 903–906 Field test, 931 Finite element, 94, 340, 1067–1069, 1074, 1185, 1499, 1503, 1624, 1693 Finite element model, 51, 52, 60, 93, 102, 322, 339 Flexural bending strength at break against untreated samples, 88 Flexural bending strength at yield against treated samples, 87 Fluidized gas at the bottom and ratio of the bed height to diameter, 65 Fluxed pellets, 17, 18, 20 Forced convection, 1007, 1008, 1012, 1015, 1016, 1184, 1441–1443 Four acids microwave-assisted digestion and chemical analysis by ICP-OES, 839 Fractal dimension, 979, 980, 984, 991 Fractography, 451 Fracture mechanics, 1090, 1091 Fracture toughness, 408 Freeze lining, 1149, 1155, 1157, 1158, 1161, 1162, 1167–1178, 1181–1194 Freeze lining in refractory based systems, 1155 Freeze-lining formation from different slag clusters, 1173 Furnace, 5, 17, 19, 41, 63–66, 69–72, 74, 75, 106–108, 111, 145, 188, 222, 423, 466, 497, 521, 522, 530, 550, 563, 572, 581, 582, 601, 630, 634, 647, 721, 729, 785, 810, 847, 856, 867, 916, 960, 1045–1048, 1152, 1188, 1067, 1072, 1080–1084, 1112, 1140, 1141, 1149–1153, 1155–1158, 1162, 1163, 1167, 1173, 1177, 1178, 1181, 1182, 1184–1193, 1200, 1218, 1227–1229, 1231, 1232, 1264, 1315, 1316, 1428, 1439, 1474, 1475, 1535, 1666, 1667 Furnace cooling, 1168, 1169, 1191 Furnace lining development within ISASMELT™, 1157 Future directions, 1658 G Galfan coating, 949–953, 955, 956 c!d Phase Transformation in DSS, 1451 Gas-insulated switchgear, 137 Gas-sensitive response to low concentrations of three sulfides, 139 Gas–solid flow distribution characteristics, 63 Gears, 1589, 1590, 1593, 1594, 1597–1599, 1622 Gemba Walk for Lean Waste (Muda), 1082

1721 Geometrically Necessary Dislocations (GNDs), 538, 542, 543, 545 Geometric structure of grains, the, 983 Geometry and mesh, 32 Geopolymer, 77 GH4169 alloy, 347, 348 Glassy alloys, 879, 883 GLMA conferences Grade components, 1297 Grain boundary nucleation, 1000 Grain growth, 225, 332, 347, 348, 399, 521, 632, 866, 887, 888, 897, 921, 979–986, 988, 989, 991, 1238, 1429, 1443, 1665, 1666, 1668, 1671, 1672, 1674 Grain growth and effective parameters, 983 Grain refinement, 280, 486, 487, 490, 492, 529, 530, 533, 534, 551, 556, 588, 638, 809, 813, 816, 865, 1259, 1384, 1441, 1666, 1693 Granularity distribution, 21 Graphene, 173–182, 681–683 Graphite, 145, 175, 218, 222, 224–229, 231, 232, 234, 476, 682, 734, 781–783, 787, 810, 831, 893, 901–906, 916, 919, 950, 1141, 1167, 1169, 1170, 1200, 1264, 1314, 1374, 1379, 1392, 1666, 1702 Growth of the light metals alliance, 1690 H Hardness, 78, 144, 205, 209, 213, 215, 230, 234, 239, 241, 243, 244, 247, 251, 256, 261, 262, 264–266, 269, 271–277, 281, 285, 496, 499, 524, 579, 599, 601–605, 630, 631, 634, 699, 703, 711–713, 715, 717, 733–740, 742, 760, 761, 809, 810, 813, 816, 868, 871, 872, 887, 916, 928–930, 1099, 1101, 1103, 1105–1108, 1253, 1254, 1263, 1265, 1266, 1268, 1272, 1292, 1294, 1328, 1329, 1331, 1421–1425, 1499, 1502–1505, 1534, 1578, 1622, 1623, 1627 Hardness evolution, 1253 Hardness testing, 813 Hardystonite, 143–151, 153, 154 Haynes 230, 282, 505, 1313–1315 HCB interfacial microstructure of pure iron, 6 HCl acid, 927, 932–936 Heat-Affected Zone (HAZ), 214, 215, 231, 234, 241, 242, 245, 734, 737, 742, 1089, 1101, 1103, 1105, 1107 Heat transfer, 52, 55, 173, 180, 182, 194, 319, 320, 347, 879, 880, 885, 1068, 1107, 1161, 1162, 1167–1169, 1172, 1173,

1722 1175, 1177, 1178, 1182–1187, 1190, 1191, 1193, 1283, 1629, 1704 Heat transfer model, 52 Heat treatment, 5–7, 154, 217–219, 221, 222, 226, 252, 377, 384, 387–390, 392, 398–400, 413–415, 417, 428, 437, 444, 456, 464–466, 490, 504, 508, 522, 530, 574, 580, 688, 720, 721, 724, 728, 729, 773, 959, 960, 965, 980, 1100, 1108, 1268, 1298, 1379, 1384, 1385, 1402, 1421, 1423, 1425, 1438, 1656 Heterogeneous nucleation, 1063, 1064, 1441, 1443, 1665–1667, 1674 High Cycle Fatigue (HCF), 224, 410 High cycle fatigue performance, 224, 232 High-Entropy Alloys (HEAs), 1263, 1271, 1281, 1282, 1284 High strength and high conductivity, 629, 630, 633, 634, 1373, 1374, 1380 High temperature, 17, 24, 106, 111, 213, 255, 456, 497, 498, 521, 621, 637, 638, 645, 646, 682, 767, 781–783, 787, 820, 847, 916, 928, 960, 980, 999, 1038, 1047, 1058, 1061–1064, 1068, 1069, 1072, 1140, 1144, 1147, 1156, 1172, 1200, 1264, 1294, 1313, 1314, 1358, 1384, 1427, 1447, 1448, 1472, 1483, 1492, 1495, 1535, 1631, 1655 High-T precipitation, 1654 Hole expansion ratio, 571, 576 Hollow, 33, 155, 156, 160–162, 165, 170, 486, 765, 766, 929, 1037, 1039–1042, 1564, 1570, 1645, 1651, 1652 Hollow fiber approach, 1651 Homogeneous nucleation, 999 Hot cells, 1325–1327, 1331, 1332 Hot compressive bonding, 3, 4 Hot-dipping, 949 Hot stage microscope, 1139–1144, 1147 Hot-top Pulsed Magneto Oscillation (HPMO), 1437, 1438, 1440, 1441, 1443 Hydrochloric acid leaching, 1227, 1228, 1233 Hydrogen-induced ductility loss, 1111, 1112, 1119–1121 Hydrogen-induced mechanical degradation, 1119 Hydrogen permeation behavior analysis, 1117 Hydrogen quench continuous annealing, 522 I Identification of significant differences in slag practice regimes, 1162 Identification of temperature dependent properties, 1169

Subject Index Identification of the clusters, 1163 Identification of the heat transfer coefficients, 1172 Impact and analysis, 1299 Impact energy results, 583 Impact of atomic radius on grain growth, 985 Impact of initial seed number on grain growth, 985 Impact of seed activation coefficient on grain growth, 985 Impact of sintering temperature on grain development, 986 Impact of sintering time on grain growth, 986 Impact toughness, 217, 231, 232, 234, 235, 421, 422, 424, 427–429, 558, 583, 1101 Inclusion modification, 561 Incoherent boundaries, 1447, 1456 Incremental Step Tests (IST), 386, 388 Induced (Non-autonomous) self-healing, 1645 Induced self-healing, 1645, 1651 Inductively Coupled Plasma Optical Emission Spectrometer (ICP-OES), 434, 811, 812, 837–841, 843, 1313, 1314, 1316, 1319, 1320 Influence of different feeding positions, 70 Influence of length of the sample, 1702 Influence of liquid–solid ratio on chromium leaching from composite cementitious material, 1221 Influence of particle size on chromium leaching from composite cementitious material, 1220 Influence of pulling speed, 689 Influence of reaction time, the, 1232 Influence of Ta on as-cast microstructures, 1059 Influence of Ta on microstructures aged at high temperature, 1061 Influence of the melting zone speed, 1700 Influence of top gun position, 72 Influence of zone pass, 1702 Influences of the current densities on the coating properties, 921 Influences of the current waveform on the coating properties, 920 Influences of the deposition parameters on the coating properties, 920 Influencing emissivity of polymeric materials, 1631 Inhibition, 434, 817, 939–944, 947, 1460, 1461 Inhomogeneity, 408, 538, 567, 1500, 1502, 1505, 1579 Initial standing wave study, 1525 Inkjet printing/material jetting, 1567

Subject Index In situ biomechanical testing coupled with micro-CT, 795 In-situ composite, 1373, 1374, 1378, 1380 Installation structure, 1067–1069, 1072, 1074 Integrated Computational Materials Engineering (ICME), 1291, 1292, 1294, 1299 Interaction effects, 1175 Intercritical annealing, 571–576 Interface characteristics, 261, 266 Interfacial oxides, 4, 5, 9, 10, 13 Interfacial oxides dissolution by AIMD, 9 Intermetallic, 239, 259, 445, 448, 456, 464, 620, 637, 643, 646, 871, 891, 960, 961, 1101, 1204, 1266, 1401–1403, 1405–1407, 1422, 1491 Intermetallic compounds, 240, 243, 247, 259–261, 263–266 Intermetallic phases, 212, 215, 260, 443–445, 452, 476, 479, 481, 959, 1284, 1286, 1401, 1483, 1492, 1494, 1495 Intrusions, 1125–1128, 1130, 1132, 1133 Iron-based alloys, 475, 476, 1058 Irradiated materials, 1325, 1326, 1331, 1332, 1342 ISASMELT™, 1149–1158 ISASMELT™ Concept, 1151 ISASMELT™ Reaction Mechanisms, 1156 Isothermal heat treatment, 222 Isothermal melt solidification, 1667 K Kaizen Event for Melt Furnace, 1082 Kinetic exponent of grain growth evolution simulated in atomic scale, 989 Kinetic model, 1045, 1046, 1048, 1299 Kinetics, 889, 988, 989, 991, 999, 1021, 1024, 1046, 1052, 1053, 1125, 1126, 1130, 1131, 1133, 1390, 1402, 1403, 1405, 1407, 1448, 1460, 1464, 1534, 1537, 1575, 1584, 1631, 1637, 1665–1667, 1674 Kinetics calculation of AlN precipitation, 999 Kurz–Fisher (KF), 1357–1368 L Laser-additive manufacturing, 206 Laser beam melting, 485, 486, 491, 492 Laser deposition, 206–208, 211, 212, 214, 215 Laser deposition on substrates with grooves, 211 Laser Engineered Net Shaping, 207, 279, 281–284, 821, 1000 Laser hardening, 733–738, 740, 742

1723 Laser heat source, 321 Laser powder bed fusion, 395, 504, 505 Lath martensite, 457, 459, 529–534, 575 Layered martensite, 584, 588, 589 Layer thickness and layer characterization, 724 Leaching risk, 1217, 1223 Lead, 1150, 1153 Lead free, 645, 646 Level set, 319, 323, 324 Light metals, 240, 332, 1687–1692 Light metals technology, 1690 Limitations of contactless gears, 1593 Local ionic structure information, 1361 Low-Cycle Fatigue (LCF), 401, 410 Low-T precipitation, 1654 M Magnesium, 4, 17–25, 422, 475, 809, 841, 843, 856, 950, 1046, 1090, 1199, 1200, 1228, 1229, 1232–1234, 1306, 1385, 1484, 1688, 1689 Magnesium sheet, 1693 Magnetic Augmented Rotation System (MARS), 1589–1593, 1597–1599 Magnetic separation, 839, 840 MAM kinematics, 1500 Manufacturing parameters and post-processing, 506 Maraging steels, 455, 456 Martensitic phase, 448 Martensitic phase transformation, 433, 434 Material characterization, 281 Material safety and the environment, 1571 Materials for microneedles, 1570 Materials preparation and characterization, 867 Mathematical model, 1070 Measurement issues, 1083 Mechanical activation, 1533–1535, 1537, 1539, 1540 Mechanical behavior, 396, 414, 496, 499, 507, 520, 579, 646, 1272, 1306, 1309, 1325, 1326, 1348, 1390 Mechanical behavior analysis, 499 Mechanical property, 77, 78, 84, 89, 143, 144, 174, 180, 217–219, 232, 252, 253, 261, 262, 264, 266, 269, 270, 279, 280, 332, 347, 355, 356, 395, 398, 403, 407, 414, 422, 424, 433–436, 443, 444, 450, 451, 465, 469, 475, 476, 486, 487, 495–497, 499, 500, 510, 514, 519–522, 525, 526, 538, 550, 554, 571–574, 579, 581–583, 589, 629, 634, 637, 638, 645, 647, 682, 711, 714, 715, 734, 793, 810, 817, 820, 845, 846, 852, 853, 855, 856, 861–863,

1724 884, 902, 959, 1100, 1101, 1200, 1209, 1237, 1263–1265, 1269, 1271, 1281, 1298, 1305, 1306, 1309, 1335, 1348, 1373, 1379, 1390, 1402, 1423, 1428, 1456, 1483, 1484, 1493–1495, 1575, 1579, 1581, 1586, 1643, 1646, 1691 Mechanical testing of structural materials for the NPP’s, 1327 Mechanism, 4, 6, 9, 12, 13, 18, 170, 203, 232, 233, 349, 358, 395, 422, 469, 513, 529, 537, 562, 698, 700, 701, 705, 706, 708, 709, 716, 746, 747, 814, 908, 915–917, 924, 939, 943, 1022, 1037, 1038, 1042, 1119, 1125, 1130, 1167, 1174, 1200, 1203, 1257, 1258, 1274, 1314, 1337, 1352, 1353, 1358, 1377–1379, 1392, 1412, 1443, 1448, 1455, 1483, 1484, 1487, 1492, 1495, 1530, 1533, 1539, 1540, 1646, 1656, 1659, 1669, 1674 Medical Mg alloys, 1484 Medium-carbon steel, 1389, 1399 Melting, 23, 198, 200, 210, 226, 240, 247, 259, 260, 262, 279, 292, 295, 319, 325, 339–343, 345, 408, 415, 422, 423, 438, 467, 476, 485, 487, 490, 503, 599, 634, 637, 646, 647, 734, 735, 740, 746, 748, 756, 781, 783, 810, 811, 828, 852, 866, 1007, 1029, 1139–1141, 1146, 1147, 1168–1172, 1174, 1175, 1177, 1178, 1218, 1264, 1267, 1313, 1314, 1374, 1379, 1403, 1428, 1438, 1442, 1447, 1448, 1450, 1452, 1456, 1488, 1566, 1570, 1655, 1666, 1697–1704 Melting behaviour, 1139, 1140 Melt loss, 1079–1081 Melt pool, 198, 269–277, 319, 320, 322, 323, 325–328, 339, 340, 342–345, 399, 408, 458, 464, 486 Metakaolin, 77 Metallographic analysis, 740, 1103 Metallographic and micro-structural analysis, 736 Metallographic investigations, 389 Metallography, 222, 601 Metal powders, 187, 355, 867 Methodology, 414 Methods of measuring burn rate, 1637 Methods of measuring emissivity, 1635 MgCl2-KCl, 1313, 1314, 1316–1320 MgO, 4, 17–19, 22–25, 45, 828, 1045, 1046, 1048, 1052, 1053, 1140–1144, 1155, 1163, 1166, 1170, 1177, 1218, 1219, 1229–1233 Micro-ECM, 746

Subject Index Micro-EDM, 745, 746 Micro-hardness, 734, 736–739 Micro-hardness analysis using Taguchi experimental design, 738 Micro-structure, 733–736, 739, 740, 742 Micro-structure and properties, 630, 632 Micro-structure and properties of Cu–0.3%Ag alloy, 630 Microalloyed steel, 537–540, 542, 543, 545, 546, 1237 Microcapsulation approach, 1652 Micrograph analysis, 623 Microhardness, 209, 214, 219, 223, 230, 232, 234, 241, 243, 244, 247, 251, 252, 255, 257, 279–281, 283, 284, 630, 811, 813, 814, 867, 871, 873, 1105, 1265, 1266, 1422, 1483, 1486, 1493–1495 Microhardness and Young’s Modulus, 871 Microhardness measurements, 1105 Microhardness study, 243 Micropillar, 1306, 1307 Microscopy study, 242 Microsphere, 159, 160, 167 Microstructural analysis, 262, 479, 582 Microstructural characterization, 812 Microstructural characterization after the SCC experiment, 1094 Microstructural characterization of As-received sample, 1093 Microstructural evolution, 283, 583 Microstructural path, 1464 Microstructure, 5, 6, 8, 18, 20, 143, 147, 150, 166–168, 186, 205, 209, 213–215, 218, 219, 224–226, 234, 251–256, 263, 266, 269, 273, 281, 283, 320, 331–334, 337, 347, 349, 356, 371–373, 375–377, 381, 391, 395, 396, 399–402, 407, 410, 421–423, 425, 426, 429, 433–435, 443–445, 448–450, 452, 455–459, 461, 464–470, 476, 478, 479, 481, 486, 490, 495, 496, 505, 513, 520, 521, 524, 525, 531, 537, 538, 545, 546, 550, 551, 553, 556–559, 571, 572, 574–576, 580, 584, 585, 600, 601, 603, 604, 620, 637–643, 646, 712, 767, 769, 799, 800, 806, 807, 809–812, 814, 849, 851, 853, 866, 867, 902, 903, 908, 909, 911, 959, 979, 980, 984, 986, 988, 991, 1007, 1063, 1093, 1100, 1101, 1103, 1104, 1107, 1111, 1113, 1114, 1116, 1118, 1121, 1183, 1187, 1191, 1204, 1239, 1256, 1258, 1264, 1265, 1267–1269, 1272, 1292, 1294, 1373–1376, 1379, 1380, 1383–1385, 1390, 1391, 1402, 1403,

Subject Index 1405, 1415, 1418, 1427–1430, 1432, 1437–1439, 1441, 1447, 1448, 1461–1466, 1483–1486, 1491, 1494, 1495, 1648, 1649, 1655, 1665 Microstructure analysis, 551 Microstructure and microhardness, 251, 253, 256, 279 Microstructure evolution, 332, 402, 558, 571–575, 991, 1062, 1063, 1254, 1274, 1306, 1401–1403, 1484, 1486 Microstructure of the WAAM-NAB alloy and comparison with cast NAB, 445 Microstructure-Oriented Scientific Analysis of Irradiated Concrete (MOSAIC), 509, 902, 907, 911 Microtexture, 463, 465, 470 Mid-week begins the laboratory portion of the class, 1296 Mild steel, 240, 697, 699, 700, 704, 709, 939, 940, 947 Milling time, 849, 865, 869–873, 1537 Minerals, 816, 908, 911, 1038, 1140, 1223, 1224, 1228, 1522 ML40Cr, 997, 999, 1004 Modeling, 30, 289, 292, 319, 320, 322, 323, 325, 348, 355, 356, 521, 1007, 1019, 1291–1294, 1298, 1299, 1412, 1673 Modeling technique, 322 Model validation, 54 Model verification, 33 Modulation-assisted machining, 1499, 1500, 1505 Molecular dynamics, 765–769, 771, 772, 884, 1272, 1348, 1357–1360, 1368 Molecular dynamics simulations, 767, 1271, 1305, 1347, 1361 Molecular interdiffusion, 1650 Molecular oxygen adsorption and dissociation on the (001) surface, 1041 Molten salt, 902, 915–919, 924, 925, 1200, 1314, 1357–1368 Monodispersed, 156, 160 Monte carlo model, 982 Monte carlo simulation, 126, 130, 132 Morphology of rapid quenching SmFe ribbons, 1434 Motivation, 253, 290 Mount Isa Mines copper plant, 1153 Moving boundary, 879, 1019, 1023 Mr. Novikov model, 982 Multicomponent, 366, 1007, 1008, 1010, 1015, 1019, 1021, 1023–1025, 1281 Multi-dendrite growth, 1012

1725 Multiphase, 30, 266, 519, 520, 571, 574, 1008, 1019, 1023–1025, 1402 Multiphase-field, 1007, 1008, 1015 Multiphase-field model, 1008 MXRF and EDS elemental analysis results and discussion, 909 N NaCl, 492, 564–567, 681, 683, 748, 749, 755, 757, 758, 869, 915–917, 919, 922, 924, 925, 951, 959–965, 968, 969, 1090, 1200, 1253, 1255, 1257 Nacre, 1648 Nanomaterial, 174 Nanoparticles, 145, 146, 185–194, 485, 486, 488, 492, 1347, 1348, 1350, 1655 Nanosheets, 165–168, 170–172, 681, 682 NanoSMA-dispersoids, 1655 Natural composites, 1646 Natural fibre, 78 Nb rubric, 1297 NdFeB, 1590 Neutron diffraction, 513, 514, 516 Neutron irradiation, 901–905, 907, 908, 1335 Nickel, 167, 187, 205, 212, 215, 218, 251, 259–263, 265, 269–271, 277, 320, 347, 352, 414, 434, 444, 448, 464, 466, 504, 505, 841, 843, 866, 870, 959, 1007, 1010, 1015, 1037, 1038, 1040–1042, 1058, 1107, 1150, 1186, 1188, 1314, 1412, 1595 Nickel Aluminum Bronze (NAB), 205–207, 212, 213, 215, 443–452 Nickel-Titanium, 1412 Nickle-based superalloys, 504 Niobium, 464, 505, 712, 782 Ni–Ti Alloy, 413, 414 Nitrogen doping, 138, 436 Non-destructive inspection, 1102 Non-Gaussian beams, 321, 326 Non-standard techniques SCC with X-ray computed tomography, 1329 Normal operating conditions and heat loss, 1182 Notebook rubric, 1298 Nuclear power plants, 907, 908, 1326, 1335 Nucleation, 422, 426, 468, 485, 491, 492, 550, 562, 586, 716, 751, 773, 870, 883, 915, 925, 961, 997–1004, 1019, 1024, 1025, 1063, 1091, 1240, 1241, 1259, 1269, 1271, 1274, 1277, 1336, 1337, 1348, 1384, 1385, 1390, 1397, 1398, 1405, 1407, 1423, 1429, 1430, 1434, 1436,

1726 1441–1443, 1448, 1450, 1451, 1453, 1454, 1456, 1459–1461, 1464, 1465, 1467, 1654, 1665–1674 Nucleation mechanism, 997, 999, 1002, 1003 Nucleation mechanism of AlN in austenite, 999 Numerical set up and calibration curves, 1624 Numerical simulation, 347, 392, 766, 1068, 1402, 1499, 1503, 1622 O Obsolete smartphones, 838, 840 OM/SEM/TEM analysis of microstructure, 1374 Open Circuit Potential (OCP) measurement, 963 Optical microscopy analysis, 961 Optimization of grain growth in the atomic scale, 984 Optimization of media injection angle, 1608 Orthogonal arrays, 207 Ostwald ripening, 1024, 1579, 1583, 1586 Oxide layers, 719–721, 729 P Packing density, 355–358, 365, 366 Packing density for ternary powder mixtures, 358 Pair–correlation function, 1462 Parameter optimization, 202 Particles powder morphology, 869 Percent composition, 361 Periodontal ligament, 797 Perspective and course evolution, 1291 Pharmaceuticals, 290, 1563 Phase diagram, 6, 23, 263, 264, 449, 477, 479, 631, 646, 647, 812, 896, 1028, 1032–1034, 1140, 1141, 1143, 1390, 1472, 1485, 1490–1492 Phase-field, 29–31, 33, 1008, 1010, 1011, 1014, 1401–1403, 1405–1407 Phase identification, 1266 Phase transformation, 144, 213, 280, 434, 445, 520, 549, 556, 558, 559, 600, 1019, 1238, 1263–1269, 1390, 1402, 1447–1451, 1453, 1454, 1456, 1491, 1655 Phase transition, 765, 773, 776, 1008, 1268, 1448, 1657 Photo-induced healing, 1651 Photopolymerization, 1568 Photopolymerization of acrylate resins, 1578 Physical model, 63, 1069, 1070 Physical properties, 1376 Pickling and cold rolling, 522

Subject Index Pin on disc, 711, 713 Pitting, 152, 562, 725, 750, 751, 816, 817, 923, 933, 934, 959, 961, 1251, 1259, 1317 Plasma modification, 819–821, 823, 825 Plasticity, 177, 180, 182, 377, 513, 519, 550, 915, 1112, 1119, 1120, 1272, 1306, 1307, 1347, 1348, 1352 Point processes, 1460, 1461, 1463, 1464, 1468 Polarization curves, 145, 151, 564–566, 922, 924, 940, 951, 953–955, 967–973, 975, 1199, 1202, 1203, 1255, 1257 Polylactic acid, 292, 1207 Polymeric, 77, 156, 820, 1580, 1586, 1645 Polymeric materials, 1580, 1631, 1649 Polymers, 1649 Polyol method, 187 Polystyrene, 155–160, 1651 Pore characteristics, 855, 857, 863 Pore structure and porosity, 498 Porosity, 23–25, 51, 52, 55, 74, 93, 94, 148, 202, 213, 239, 240, 242, 244–247, 279, 323, 333, 342, 357, 395, 396, 398, 403, 409, 410, 444, 445, 452, 456, 485–487, 491, 495–500, 503, 505, 507, 508, 687, 689, 691–693, 695, 845–847, 852, 853, 855, 858, 859, 861–863, 865–868, 870, 873, 880, 901–906, 980, 984, 986, 988, 1100, 1103, 1171 Porosity and density, 870 Porous Ti, 495–500, 846, 848, 852, 853 Post-Weld Heat Treatment (PWHT), 217, 221, 233, 1099, 1100, 1107 Potential drop, 170, 1089, 1091–1093, 1095 Potential for chemical wear at constant temperature, 1177 Potential parameters, 1358 Potentiodynamic polarization and EIS study, 1255 Potentiodynamic polarization studies, 942 Potentiodynamic polarization test, 961 Powder bed fusion, 331, 504, 505 Powder consolidation, 887–889, 892, 897 Powder feeding, 206 Powder material, 506 Powder metallurgy, 355, 434, 495–497, 500, 504, 845, 855, 856, 863, 866, 873, 888, 1264, 1484 Powder surface inoculation, 488 PRDFs and CN, 1362 Precipitates and microstructure analysis, 1114 Precipitation, 40, 45, 48, 144, 214, 280, 347, 426, 443, 444, 448, 456, 459, 490, 504, 522, 537, 539, 540, 553, 561, 580, 586, 588, 725, 729, 730, 813, 816, 943, 947,

Subject Index 961, 997, 999, 1002, 1003, 1019, 1024, 1057, 1061, 1062, 1064, 1065, 1100, 1101, 1111, 1113, 1114, 1223, 1237, 1238, 1240, 1243–1246, 1294, 1298, 1374, 1375, 1401–1407, 1421, 1423, 1424, 1430, 1448, 1454, 1455, 1484, 1490, 1654, 1657 Precipitation strengthening, 543, 549, 550, 556, 557, 559, 1424 Precursor Powder, The, 889 Preliminary study on damaged radiant tube, 1068 Preparation and characterization of UV-curable formulations, 1579 Preparation of composite cementitious material, 1219 Preparation of composite materials, 176 Preparation of specimens, 828 Preparation of the specimens, 384 Preparation of the titanium foams, 846 Pressure, 4, 6, 18, 23, 39, 41, 44, 53, 64–69, 72–74, 97, 105–110, 125–132, 134–136, 168, 174, 176, 177, 187, 252, 260, 280, 290, 295, 322, 347, 397, 464, 497, 698, 719, 721, 722, 724, 728, 729, 755, 766, 768, 772, 774, 775, 810, 811, 821, 823, 847, 855–863, 867, 887, 888, 891, 892, 897, 928, 930, 932, 979, 981, 1027, 1028, 1030, 1032–1034, 1112, 1152, 1168–1170, 1172, 1263–1266, 1268, 1269, 1316, 1325–1327, 1329, 1331, 1335, 1353, 1360, 1385, 1424, 1442, 1473, 1485, 1567, 1605, 1607–1612, 1629, 1636, 1637, 1655, 1693 Principle and characteristics of diffusion connection technology, 767 Process, 785 Process development on flat substrates, 209 Processing of lab-based carburized material, 1529 Processing parameters, 521 Process parameters, 197, 199, 205, 207–212, 215, 241, 252, 253, 323, 333, 384, 395, 398, 456, 622, 624, 683, 688, 733–735, 740, 745, 748, 775, 776, 1034, 1182, 1183, 1187, 1190, 1191, 1193, 1566, 1618, 1619, 1693 Properties of fine wires with different drawn strains, 631 Prosthesis, 1208, 1484 Prototyping, 289, 296, 383, 828 PTFE-alumina coatings, 756, 758, 760, 761 PTT curve, 1001, 1003, 1004

1727 PTT curve of AlN precipitation in austenite, 1001, 1003 Pylon, 1207, 1209, 1210, 1213, 1214 Q Quantification of effect of different regimes on freeze-lining behaviour, 1173 Quasi-static properties, 399 R Radial distribution function, 765, 770, 1360, 1362 Radiation temperature effect, 1339 Rapid quenching, 490, 887, 1427, 1429–1432, 1434 Rapid solidification, 253, 273, 464, 1423, 1429–1431, 1436 Raw materials, 174 Raw materials for experiment, 64 Recommissioning of furnaces, 1082 Reduction amount, 97 Reduction efficiency of convex rolls, 57 Reduction factor, 361, 366 Reduction force, 97 Reduction position, 99 Relative density, 398 Repeatability of gas sensor, 140 Residual stress, 252, 253, 269, 273, 396, 400, 439, 620, 929, 1239 Resistivity, 599–605, 619–621, 623, 626, 632, 746, 1374 Response characteristics to three sulfides, 139 Retained austenite, 217, 218, 223, 226–232, 234, 235, 458, 459, 461, 513, 525, 549, 550, 553, 554, 557–559, 584, 585, 1118, 1390 Ringer solution, 143, 145, 151 Rolling, 241, 252, 422–425, 428, 514, 521, 522, 526, 530, 564, 572, 573, 581, 583, 586, 588, 599–605, 632, 809–812, 930–932, 1112, 1298, 1385, 1483, 1487, 1488, 1492, 1495, 1604 Roll surface profile, 93, 94, 102 Rolls used in soft reduction, 94 S Seebeck effect, 781, 782 Segregation, 51–53, 55, 59, 60, 93, 97, 227, 247, 449, 464, 534, 866, 959, 1012, 1100, 1118, 1119, 1165, 1238, 1244, 1427–1429, 1436, 1448, 1668, 1669, 1674, 1698–1700, 1704 Selective Laser Melting (SLM), 320, 332, 339–341, 343–345, 371, 374, 383, 384,

1728 395, 396, 399, 400, 407–410, 433–440, 476, 482, 491 Self-diffusion coefficient, 1366 Self-healing, 1643–1646, 1648–1659 Self-healing concept, 1643 Self-healing concrete and asphalt, 1658 Self-healing via reversible bond formation, 1651 Self-propagatinghigh-temperature synthesis, 1281, 1282 SEM analysis, 705 SEM/EDS analysis, 1105 Semi-coherent boundaries, 1447, 1450, 1456 Semi-solid powder injection moulding, 1483–1487, 1492, 1495 Senior design, 1292–1294, 1299, 1300 Sensors, 140, 186 Sequential images of the grain growth evolution, 986 Sequential machining, 746 Severe plastic deformation, 1375, 1418 Shape Memory Alloys (SMAs), 413, 638, 1411, 1412 Sharp interface, 29, 1019, 1021 Shear-punch test, 800 Shear strength, 799–801, 803–807 Shear test, 645–652 Shockwave, 887–889, 893, 897 Shockwave powder consolidation, 888 Short crack growth, 1130 Shrinkage and reduction efficiency, 97 SiB3, 1472, 1476, 1477 SiB6, 1472, 1476, 1477 SiBn, 1472, 1476, 1477 Si–B phase diagram, 1472 Sigma phase, 1019, 1024, 1025, 1268 Similar Metal Weld (SMW), 1099–1108 Simplified Molecular Interaction Volume Model (SMIVM), 1027–1032, 1034 Simulation cases and DIAT shift concept, 1337 Simulation details, 1348 Simulation progress, 348 Simulation under different gun positions, 72 Simulation, 4, 5, 30, 37, 55, 72–74, 96, 121, 122, 125, 126, 132, 140, 319, 320, 323, 347–350, 352, 389, 490, 510, 713, 717, 765–768, 772, 774, 979–984, 986, 989, 991, 1007, 1008, 1010–1012, 1014, 1024, 1171, 1173, 1204, 1272, 1273, 1283, 1293, 1306–1308, 1336, 1337, 1340, 1348, 1349, 1357–1361, 1368, 1402–1405, 1407, 1461, 1503, 1624, 1665, 1693, 1704 Single dendrite growth, 1011

Subject Index SiO2, 19, 22, 40, 45, 48, 562, 935, 1045–1048, 1050, 1139, 1140, 1144–1147, 1165, 1166, 1170, 1177, 1229, 1230, 1534, 1536 Six sigma team projects, 1080 Size distribution, 21, 157, 169, 355, 356, 366, 434, 465, 564, 566, 702, 848, 863, 889, 890, 1113–1115, 1223, 1238, 1463–1465, 1667 Slag, 24, 29, 30, 33, 34, 37, 64, 69, 1045–1053, 1139–1142, 1144, 1146, 1147, 1149–1151, 1155–1158, 1161–1173, 1175–1178, 1182–1184, 1186–1189, 1191–1193, 1217–1221, 1224, 1227–1234, 1283–1285 Slip markings, 1125, 1126 Slurry, 167, 697–701, 703, 705, 708, 709, 1228, 1523, 1524, 1528–1530 Slurry pipeline, 699–701 Slurry pot, 697–699, 701, 704, 705, 708 Slurry pot tester, 701 SMA-clamp and melt, 1655 Sm content, 1427, 1431, 1434, 1436 Smelting, 17, 105–111, 1028, 1149–1153, 1155, 1157, 1170, 1182, 1193, 1218, 1220, 1227, 1699, 1700, 1702 Sm-Fe alloy, 1428 Sodium dodecyl sulfate, 155, 157, 162 Soft reduction, 51–53, 55, 57–60, 93, 94, 96 Soft reduction amount, 59 Solder, 637, 638, 645–652, 1655–1657 Solder tubes/capsules, 1655 Sol-gel, 143, 145, 146, 153, 687, 688, 690, 691, 694, 1533, 1534 Solid particle used, 702 Solidification, 51–57, 59, 60, 95, 96, 99, 212, 214, 225, 242, 247, 273, 277, 319, 331, 334, 335, 337, 339–341, 343–345, 347–352, 407, 422, 445, 448, 464, 467, 468, 485–487, 490, 504, 521, 580, 602, 637–642, 749, 769, 870, 879, 1007, 1008, 1010–1016, 1058, 1059, 1170–1172, 1183, 1223, 1242, 1428–1430, 1436–1439, 1441–1443, 1447, 1448, 1450, 1452, 1456, 1486, 1490–1492, 1657, 1665–1674, 1690, 1698, 1699 Solidification shrinkage, 55 Solidified microstructure, 1437, 1438, 1440, 1441, 1443 Solid-state electrochemical cell, 1474 Solute carbon, 529, 530, 534 Solute-depleted regions, 1237–1239, 1242–1246

Subject Index Solute distribution of slab in solidification, 1242 Solute-rich regions, 1237, 1242, 1245, 1246 Spark plasma sintering, 866, 888, 1263, 1264 Specific surface area and pore structure distribution, 21 Specimen preparation, 794 Stainless steel, 4, 105–108, 111, 143–145, 151, 153, 154, 198, 199, 218, 219, 251, 253, 259–263, 266, 279, 280, 409, 456, 562, 687, 688, 698–700, 785, 802, 828, 928, 930, 960, 1019, 1024, 1090, 1099, 1107, 1131, 1161, 1162, 1200, 1217–1220, 1223, 1224, 1258, 1329, 1438, 1447, 1570 Standard techniques, 1327 Standing wave, 1521, 1523–1530 Standing wave processing, 1523 Statistical analysis, 831 Statistical simulation research method, 982 Statistics of structure information and self-diffusion coefficient, 1360 Steady-state freeze lining thickness, 1185 Steel, 17, 18, 29, 30, 33, 34, 37, 51–54, 56, 93, 94, 96, 102, 105–111, 148, 149, 152, 153, 225, 239–245, 347, 349, 421–429, 455, 461, 486, 506, 513–521, 525, 529, 530, 537, 539, 540, 542, 543, 545, 546, 550–555, 557, 558, 561–568, 571–576, 580, 582–589, 599, 600, 603, 621, 639, 698, 699, 702, 713, 717, 720–722, 725, 729, 730, 735, 746, 747, 781–783, 787, 821, 856, 867, 887, 891, 908, 915–917, 919–922, 924, 925, 928, 939, 940, 943, 947, 949, 950, 953–955, 959, 960, 967–969, 972–975, 997–999, 1004, 1024, 1085, 1099–1101, 1106, 1107, 1111–1121, 1131, 1132, 1162, 1167, 1169, 1170, 1174, 1176, 1181, 1184, 1200, 1201, 1204, 1205, 1218, 1228, 1237, 1242, 1245, 1246, 1251–1254, 1258, 1259, 1325–1332, 1336, 1389–1392, 1394, 1395, 1398, 1401–1406, 1437, 1439, 1440, 1447–1449, 1472, 1567, 1606, 1619, 1620, 1622, 1625–1627 Steelmaking, casting and hot rolling, 521 Sterolithography, 1575 Strain concentration, 793, 796, 797 Strain during cutting, 1501 Strain glass, 1411, 1412 Strain glass alloys, 1411 Straining temperature effect, 1340 Strain path change, 542, 546

1729 Strength, 4, 17–19, 22–25, 77–79, 82–84, 87–89, 96, 122, 144, 154, 173, 177, 179, 180, 182, 199, 201, 202, 217, 218, 232, 233, 235, 240, 259, 260, 262, 266, 280, 283, 292, 332, 336, 337, 342, 347, 371, 376, 396, 399, 400, 402, 408–410, 421, 422, 433–437, 440, 443, 451, 452, 455, 456, 464, 475, 485, 486, 492, 495–498, 500, 504, 505, 513, 515, 516, 519, 520, 524, 525, 529, 531–533, 537, 549–551, 553, 556, 557, 561, 562, 571–576, 579–581, 583, 586, 589, 629, 633, 638, 639, 643, 700, 712, 720, 755, 800, 810, 845, 846, 848, 853, 855, 856, 862, 880, 903, 908, 928, 930, 950, 959, 964, 980, 997, 1058, 1070, 1101, 1112, 1119, 1120, 1199, 1207, 1208, 1210, 1211, 1214, 1218, 1263, 1264, 1272, 1294, 1305, 1308–1310, 1338, 1348, 1373, 1379, 1384, 1390, 1401–1404, 1406, 1407, 1416, 1421–1424, 1428, 1483, 1484, 1495, 1533, 1534, 1570, 1590, 1598, 1606, 1618, 1630, 1647, 1648, 1650, 1652, 1657 Stress and strain distribution, 99 Stress Corrosion Cracking (SCC), 1089–1091, 1094–1096, 1329 Stress raiser, 1102 Stress–strain system, 375 Structure element, 371, 376, 378 Submerged arc, 1162, 1184, 1186 Superalloy, 347, 349, 352, 463–465, 503, 1007, 1008, 1010–1012, 1015, 1460 Surface analysis, 722, 727 Surface analysis before and after wear testing, 727 Surface inoculation, 487, 488, 492 Surface integrity, 279, 280, 745–747, 750–752, 1620 Surface morphology, 951 Surface morphology and sample preparation for SEM observation, 831 Surface morphology using micro-structure, 739 Surface oxidation, 1037 Surface properties, 1623 Surface relief and crack initiation, 1127 Surface roughness, 161, 261, 279–284, 395–398, 409, 734, 745–748, 750–752, 757, 758, 766, 830–832, 853, 1572, 1622–1624 Surfactant, 156, 158, 1521, 1523–1530 Synergetic induced precipitation behavior of c phase, 1454 Syntheses of PS microspheres, 157

1730 Synthesis of samples, 166 T Tafel analysis, 681, 683, 684 Tafel curve, 1201 Taguchi method, 734, 735 Tantalum, 198, 712, 1139, 1140, 1146, 1147 Technical considerations and regulatory challenges, 1570 Temperature, 1045–1047, 1050–1053 Temperature effect, 1231, 1338–1340, 1342 Temperature result, 325 Tempering, 522, 530, 575, 579, 583, 585, 586, 588, 589, 1046, 1111, 1112, 1114, 1115, 1117–1121 Tensile property, 82, 86, 223, 231, 240, 399, 400, 414, 450, 463–465, 469, 470, 550, 637–639, 1693 Tensile strength, 79, 82, 84, 91, 223, 262, 399, 401, 408, 409, 433, 435, 437, 440, 450, 463–465, 469, 470, 513–515, 519, 520, 524, 525, 554, 571–573, 576, 589, 629–634, 638, 642, 783, 810, 856, 903, 1070, 1072, 1090, 1100, 1329, 1374, 1378–1380, 1448, 1655 Tensile strength at break against treated samples, 87 Tensile strength at yield against untreated samples, 84, 89 Tensile stress and tensile strain, 99 Tensile tests, 582 Terminal velocity and floating shape, 34 Test setup, 374 Test specimen, 701 TG analysis, 1535 Theoretical analysis, the, 1699 Theoretical curves, 967, 969, 970, 973, 976 Theory and simulation procedure of grain growth ceramic, 981 Thermal analysis, 199 Thermal conductivity and density of the freeze-lining, 1171 Thermal damages, 745, 746, 748, 752 Thermal decomposition, 173, 177 Thermal decomposition characteristics of composite, 177 Thermal history, 251, 270, 275 Thermal-mechanical behavior, 93, 94, 102 Thermal stress, 342, 688, 689, 959, 1068–1072, 1074, 1273 Thermal stress model of radiant tube, 1069 Thermo-Calc, 1059–1062 Thermocouple, 781–783, 785–788, 960, 1172, 1173, 1474, 1475

Subject Index Thermocouple design, 783 Thermodynamic calculation, 106 Thermodynamic model, 1237, 1238, 1247 Thermodynamic properties, 1038, 1471, 1472, 1478 Thermodynamics, 565, 885, 943, 1348, 1648, 1649 Thermodynamics of Self-healing, 1648 Thermoelectric power, 782, 785 Thermomechanical processing, 588 Thermoplastic polymers, 1650 Thermoset polymers, 1651 Thickness, 4–6, 8, 13, 30, 31, 52, 54–56, 59, 96, 97, 99, 144, 145, 152, 167, 198, 240, 241, 262, 270, 281, 295, 324, 340, 342, 345, 391, 397, 423, 459, 466, 514, 521, 522, 530, 531, 563, 572, 599, 601, 603, 619–621, 623, 626, 640, 647, 683, 687, 689, 690, 694, 695, 701, 719, 721, 724, 725, 756, 766, 799–806, 810, 811, 857, 858, 879, 883, 885, 891, 917, 920–924, 929–932, 951, 972, 973, 982, 998, 1022, 1102, 1107, 1113, 1117, 1127, 1167, 1168, 1173, 1175, 1178, 1181–1183, 1185–1191, 1193, 1233, 1242, 1254, 1256, 1265, 1268, 1274, 1276, 1315, 1384, 1403, 1413, 1414, 1417, 1418, 1439, 1448, 1501–1503, 1505, 1567, 1606, 1636, 1637, 1652, 1700, 1704 Three conditions for density calculation, 362 Three-dimensional printing, 1563, 1564 Ti-6Al-4V, 395–404, 771, 827–829, 835 Ti-6Al-4V alloy, 745, 747–751 Ti6-Al4-V alloy molecular dynamics simulation, 767 TiC reinforcement, 495, 500 Ti-Ni alloy, 869, 872 TiO2, 137–142, 991, 1140–1142, 1144, 1218, 1227–1229, 1231, 1233, 1282 Titanium, 4, 138, 144, 251, 259–263, 265, 266, 414, 422, 488, 496, 530, 712, 735, 756, 765, 772, 775, 810, 827–829, 845–847, 850, 855–863, 866, 867, 870, 871, 1100, 1101, 1103, 1227–1229, 1231–1234, 1383, 1384, 1570, 1688 Titanium alloy, 766, 767, 772, 866, 873 Titanium-containing blast furnace slag, 1227, 1228, 1230, 1231, 1233 Titanium foam, 855, 856, 858, 859, 861–863 Topology simulation research method, 982 Toughness, 78, 79, 177, 179, 180, 182, 218, 233, 407–410, 421, 422, 443, 549, 550, 553, 554, 556, 557, 559, 561, 562,

Subject Index 579–581, 589, 711–713, 715, 717, 720, 800, 915, 1107, 1294, 1314, 1327, 1328, 1331, 1340, 1348, 1647, 1648, 1693 Transdermal drug delivery systems, 1563 Transient freeze lining formation, 1189 Transmission efficiency of MARS, 1597 Transmission Electron Microscopy (TEM) , 185, 189–193, 219, 240, 435, 438, 440, 444, 445, 448, 450, 455–457, 461, 490, 491, 551, 552, 571, 573, 575, 901, 903–907, 980, 985, 988, 1115, 1127, 1129, 1237–1239, 1252, 1253, 1256, 1325, 1335, 1374, 1376, 1377, 1411–1413, 1415, 1417, 1533–1535, 1539, 1540 Transverse tensile, 231 Transverse tensile test, 223 Tribological characterization, 868 Tribological tests, 717, 868 Transformation-Induced Plasticity (TRIP), 513–519, 1390 Top Submerged Lance (TSL), 540, 542, 551, 1149, 1150, 1158 Tungsten carbide, 712, 717, 1521–1527, 1529, 1530 Twin boundary, 1276, 1305, 1306, 1309, 1310 U Ultra-fine wires, 629 Ultrasonic dispersion, 173, 174 Ultrasonic micro-forging, 251–257 Ultraviolet–Visible Spectroscopy (UV/Vis), 189, 190 Undergraduate, 1291, 1292 Understanding regulatory issues, 1572 Understanding the role of Ta on microstructure evolutions at high temperature, 1063 Uniform elongation, 513–515, 517, 518, 1090 U-type sealing regulating valve of feeding loop, 69 UV curable resins, 1575, 1579 V Vapour Liquid Equilibrium (VLE), 1027, 1028, 1030, 1032–1034 Vascular grafts, 820 Visual examination, 1102 VLE calculation method, 1030

1731 Volume percentage of retained austenite and it’s carbon content, 227 VS4, 165–168, 170–172 W Wear, 218, 434, 697–700, 703–705, 707, 711–713, 715–717, 719–730, 733, 734, 755–757, 760, 761, 865–867, 869, 872, 873, 929, 1082, 1155–1157, 1158, 1172, 1174, 1176–1178, 1182, 1186, 1193, 1200, 1263, 1271, 1593, 1603, 1612, 1616, 1648 Wear and COF investigations, 722 Wear behavior, 725 Wear testing, 713, 717, 727 Weight loss and EDS analysis, 1316 Workability, 599–605 W-shaped radiant tube, 1067, 1068, 1074 X X-ray computed tomography study, 245 X-ray Computed Tomography (XCT), 198, 201–203, 239, 240, 801, 905, 1331 X-ray diffraction study, 242 X-Ray Diffraction (XRD), 190, 191 XRD analysis, 223, 812 XRD results, 1536 Y Y-based rare earth, 561–564, 566, 567 Yield point, 530–534 Yield strength, 223, 262, 400, 450, 516, 522, 525, 529, 549, 554, 559, 571, 573, 576, 579, 580, 587, 589, 630, 806, 810, 852, 853, 861–863, 1070, 1072, 1090, 1329 Yttrium, 421–423, 491, 562 Z Zinc, 144, 145, 240, 497, 719, 725–728, 809, 810, 812, 841, 843, 949–956, 1153, 1184, 1186, 1199, 1200, 1204, 1205, 1609, 1610 Zirconium, 688, 782 Zn–Al–Mg coating, 949, 952, 953, 955, 956 Zone pass, 1701, 1702, 1705 Zone refining, 1697, 1698, 1701, 1702, 1704 Zone speed, 1697, 1700, 1701, 1703, 1705 ZrO2, 687–695, 1140–1142, 1144