Physics and Engineering of Metallic Materials: Proceedings of Chinese Materials Conference 2018 [1st ed.] 978-981-13-5943-9, 978-981-13-5944-6

Table of contents :
Front Matter ....Pages i-xxxiii
Microstructures and Mechanical Properties of Extruded and Aged Mg–4Zn–2Al–2Sn–(0.6Mn) Alloy (Dongqing Zhao, Yuansheng Yang, Xiaocun Song, Yu Liu, Cuicui Sun, Jixue Zhou)....Pages 1-9
Microstructural Evolution of an Al–Zn–Mg–Cu Aluminum Alloy During an Optimized Two-Step Homogenization Treatment (Hongwei Yan, Xiwu Li, Zhihui Li, Shuhui Huang, Hongwei Liu, Lizhen Yan et al.)....Pages 11-17
Microstructural Evolution and Phase Transformation of Al–Mg–Si Alloy Containing 3% Li During Homogenization (Xiaokun Yang, Baiqing Xiong, Xiwu Li, Lizhen Yan, Zhihui Li, Yong’an Zhang et al.)....Pages 19-28
Study on Stabilization Treatment of Al–Mg Alloy 5E83-H112 (Xia Wu, Hui Huang, Shengping Wen, Xiaolan Wu, Zuoren Nie, Kunyuan Gao et al.)....Pages 29-39
Simulation and Experimental Study on Hot Forging Process of SiCp/2A14 Composite (Hansen Zheng, Zhifeng Zhang, Yuelong Bai)....Pages 41-49
Effect of Boron Addition Methods on Microstructure and Mechanical Properties of a Near-α Titanium Alloy (Yingying Liu, Lihua Chai, Xiaozhao Ma, Yapeng Cui, Ziyong Chen, Zhilei Xiang)....Pages 51-64
Effect of Ultrasonic Treatment on Microstructure and Properties of Aluminum Alloy Rod Prepared by Continuous Casting and Rolling (Nan Zhou, Shuncheng Wang, Zhimin Zhang, Yuanhui Zhai, Guangyao Zhong)....Pages 65-72
Influence of High-Temperature Compression on Microstructure and Properties of Sintered Molybdenum (Fu Wang, Zenglin Zhou, Yan Li, Zhilin Hui, Xueliang He, Xiaoying Fu)....Pages 73-84
Microstructure and Mechanical Properties of a Ti–Al–Sn–Zr–Mo–Nb–W–Si High-Temperature Titanium Alloy (Y. W. Diao, X. Y. Song, W. J. Zhang, M. Y. Zhao, W. J. Ye, S. X. Hui)....Pages 85-91
Optimization of Cold-Rolling-Stabilization Process for High Mg-Containing Al Alloy (Yueying Liang, Hui Huang, Xiaolan Wu, Shengping Wen, Kunyuan Gao, Zuoren Nie et al.)....Pages 93-104
Flow Stress Behavior and Microstructural Evolution of a High-Alloying Al–Zn–Mg–Cu Alloy (Guohui Shi, Yong’an Zhang, Xiwu Li, Shuhui Huang, Zhihui Li, Lizhen Yan et al.)....Pages 105-115
Influence of Space Environment on the Properties of Diamond/Cu Composites (Zhongnan Xie, Hong Guo, Ximin Zhang)....Pages 117-122
The Effects of Electron Beam Welding Parameters on Microstructure and Properties of GH4738 Alloy (Xinxu Li, Yong Zhang, Peihuan Li, Shaomin Lv)....Pages 123-131
Microstructure and Grain Refining Performance of High-Quality Al–5Ti–1B Master Alloy (Yuehua Kang, Shuncheng Wang, Nan Zhou, Dongfu Song)....Pages 133-142
Effect of Cu Content on Microstructure and Properties of Al–Mg–Si Alloy (Hong-Xiang Li, Shengli Guo, Peng Du, Sheng-Pu Liu)....Pages 143-151
Corrosion Behavior of Inconel 625 Alloy in Na2SO4–K2SO4 at High Temperature (Yuan-Jun Ma, Yutian Ding, Jian-Jun Liu, Yu-Bi Gao, Dong Zhang)....Pages 153-166
Effects of Extrusion Conditions on Microstructure and Age-Hardening Behaviors of Al–Zn–Mg Alloy (Y. L. Wang, H. C. Jiang, D. Zhang, L. J. Rong)....Pages 167-181
Process Optimization Design of High-Strength Ag–Cu–Ni Alloy Based on Orthogonal Experiments (Helong Hu, Haibin Li, Wenjun Yu, Yongzhen Jiao, Tingyi Dong, Baoguo Lv)....Pages 183-190
Strain Distribution and Metal Flow of Bulk Forming of Molybdenum (Yihang Yang, Ailong Zheng, Zhimin Huang, Fusheng Peng, Houan Zhang)....Pages 191-200
Cold Deformation Behavior and Mechanical Properties of Forged Pure Nickel N6 (Zexi Gao, Zhi Jia, Jinjin Ji, Dexue Liu, Yutian Ding)....Pages 201-210
Microstructure and Properties of Graphene/Copper Matrix Composites Prepared by In Situ Reduction (Xu Ran, Yutong Wang, Yong Wang)....Pages 211-219
The Effect of Aluminum Content on the Microstructures of Single-Phase γ-TiAl-Based Alloy (Yaodong Xuanyuan, Yan Long, Yinbiao Yan, Sen Yang)....Pages 221-229
Effect of Al–Er–Zr Master Alloy on Grain Refinement After Heat Treatment (Haiyue Yu, Hui Huang, Zuoren Nie, Shengping Wen, Kunyuan Gao, Wei Wang)....Pages 231-240
Effect of In Situ TiB2 Particle Content on Microstructure and Properties of Cast Al–Si Alloy (Hongda Wang, Lihua Chai, He Li)....Pages 241-248
The Influence of Different Casting Methods on the Corrosion Resistance of Mg–Gd–Y Alloys (Renju Cheng, Hong Tang, Wenjun Liu, Na Zhang, Hanwu Dong, Bin Jiang et al.)....Pages 249-256
Effect of Cu Additions and Extrusion Treatment on the Microstructure and Mechanical Properties of Mg–6Sn–1Al Alloy (Zhijian Ye, Tong Li, Gui Lou, Jianhang Yue, Xinying Teng)....Pages 257-271
Strain Rate Sensitivity of GH4720LI Alloy with Two Initial Microstructures During Hot Deformation (Zhi-Peng Wan, Tao Wang, Yu Sun, Lian-Xi Hu, Zhao Li, Peihuan Li et al.)....Pages 273-284
Microstructure and Properties of Rare-Earth B4C–Copper Composites (Yan Li, Meihui Song, Yu Zhang, Yanchun Li, Xiaochen Zhang)....Pages 285-294
Evaluation of Microstructure and Refining Effect of Al–TiB2 and Al–5Ti–1B Grain Refiners (Mengke He, Lihua Chai, Hongda Wang, Ziyong Chen, Yapeng Cui)....Pages 295-306
Effects of Graphene Content and Aging Process on Mechanical Properties and Corrosion Performance of an A356.2 Aluminum Matrix Composite (Kang Wang, Jinfeng Leng, Ran Wang, Shaochen Zhang)....Pages 307-319
Microstructure and Performance of WAAM TiB2-Reinforced Al–Si-Based Composites (Yaqi Deng, Xianfeng Li, Liang Wu, Qingfeng Yang, Yanchi Chen)....Pages 321-328
Characterization of Graphene/Cu Composites Prepared by CVD and SPS (Xudong Wang, Haiping Zhang, Jiongli Li, Zhen Cao, Yue Wu, Na Li)....Pages 329-336
Hot Deformation Behavior of the As-Cast Mg–10.13Li–2.83Zn–2.78Al–0.13Si Alloy (Shouyang Gao, Defu Li, Peng Du, Shengli Guo)....Pages 337-346
Fatigue Behavior of Thermal Barrier Coated DD6 Single Crystal Superalloy at 900 °C (Jianmin Dong, Jiarong Li, Rende Mu, He Tian)....Pages 347-356
Microstructure and Room-Temperature Fracture Toughness of Nb–Ti–Si In Situ Composite Prepared by Selective Laser Melting (Yongwang Kang, Fengwei Guo, Ming Li)....Pages 357-363
Simulation and Experimental Study on the Directional Solidification Process of a Single-Crystal Superalloy Plate Casting (Runnan Wang, Qingyan Xu, Baicheng Liu)....Pages 365-375
Influence of Mo Content on the Microstructure Stability and Stress Rupture Properties of a Single Crystal Superalloy (Z. X. Shi, S. Z. Liu, X. G. Wang, J. R. Li)....Pages 377-382
Preparation and Compressive Properties of Advanced Pore Morphology (APM) Foam Elements (Yanli Wang, Lucai Wang, Hong Xu, Qiaoyu Guo)....Pages 383-390
Effect of Temperature and Strain Amplitude on Low Cycle Fatigue Behaviour of DD11 (Yuanyuan Guo, Yunsong Zhao, Jian Zhang, Yanfei Liu, Zhenyu Yang, Jiang Hua et al.)....Pages 391-398
Oxidation Mechanism of Nb–Si-Based Ultra-High Temperature Materials (Fengwei Guo, Yongwang Kang, Chenbo Xiao, Ming Li, Meiling Wu)....Pages 399-409
Microstructure and Mechanical Properties of Aluminum Alloy with Ultra-high Strength Prepared by Spray Forming (Shuhui Huang, Baiqing Xiong, Yong’an Zhang, Zhihui Li, Xiwu Li, Hongwei Liu et al.)....Pages 411-417
Effect of Pre-aging Technology on Microstructure and Mechanical Properties of 6111 Aluminum Alloy (Hongwei Liu, Shuhui Huang, Baiqing Xiong, Yong’an Zhang, Zhihui Li, Xiwu Li et al.)....Pages 419-427
Effect of Zr Content on the Microstructure, Mechanical Properties, and Corrosion Resistance of Ti–27Nb–xZr Alloys (Ying Xu, Huanhuan Wang, Yanqing Cai, Ziyan Wei)....Pages 429-438
Preparation of NiO by Precipitation Transformation and Its Supercapacitor Performance (Xinglei Wang, Yunqing Liu, Fan Zhang, Xiuling Yan)....Pages 439-452
Study on Fabrication and Compressive Properties of Mg/Al-Ordered Structure Composites (Han Wang, Yu Fu, Mingming Su, Hai Hao)....Pages 453-461
Study on Heat Treatment Process of Tungsten-Plated Diamond (Zhen Zhang, Hong Guo, Zhongnan Xie, Ximin Zhang)....Pages 463-470
Effect of Sc Modification and Pulping Process on Semi-Solid Structure of A356 Aluminum Alloy (Yuxin Zhang, Hengbin Liao, Yong Dong, Anfu Chen, Xiaoling Fu, Zhengrong Zhang)....Pages 471-484
Densification Behavior of High-Volume-Fraction (B4C + W)-Reinforced Aluminum Matrix Composites Prepared by Hot Pressing (Tiantian Guo, Changhui Mao, Shuwang Ma, Zheng Lv, Guosong Zhang)....Pages 485-493
The Finite Element Simulation of Diffusion Bonding for TiAl/Ti2AlNb Annular Structural Component (Xiaoqiang Zhang, Bin Tang, Jinshan Li, Hongchao Kou)....Pages 495-503
Effect of Hf Content on the Stress Rupture Properties of the Second-Generation Single-Crystal Superalloy with LAB (Yipeng Zhang, Jianchao Qin, Zhaohui Huang, Renjie Cui, Kai Guan)....Pages 505-514
High-Cycle Fatigue of Mg–6Gd–3Y–0.5Zr Cast Magnesium Alloys (Dehao Meng, Peijie Li, Duanzhi Wang, Wenquan Yuan, Wencai Liu)....Pages 515-525
Superplastic Forming Properties and Instability of Magnesium Alloy Sheet (Meijuan Song, Chuanhui Huang, Wenchao Jiang, Weina Lu)....Pages 527-535
The Influence of Fused Mullite Mineralizer on Mechanical Properties of Ceramic Shell (Longpei Dong, Jiansheng Yao, Bin Shen, Feipeng Guo, Xiaoyu Li, Dingzhong Tang)....Pages 537-546
Influences of Welding Current on Joint Structure and Mechanical Property of AZ31B Magnesium Board (Yingchun Tian, Ping Gao, Jianjun Lv, Mingyi Zhang)....Pages 547-559
Surface Characterization of Nickel-Base Superalloy Powder (Wenyong Xu, Yufeng Liu, Hua Yuan, Zhou Li, Guoqing Zhang)....Pages 561-567
Microstructure and Mechanical Properties of 6005A-T5 Aluminum Alloy Welded Joints by Friction Stir Welding and Metal Inert Gas Welding (Jingxuan Liu, Jian Shen, Xiwu Li, Lizhen Yan, Hongwei Yan, Hongwei Liu et al.)....Pages 569-579
Deformation Behavior and Constitutive Model for Isothermal Compression of TC4 Alloy (Jiangwei Zhong, Pan Tao, Qingyan Xu, Baicheng Liu, Zhijun Ji)....Pages 581-591
Recent Development of Hot-Pressed-/Deformed Nd–Fe–B Permanent Magnets (Ke Lv, Rui Dong, Mingjing Zhao, Guozheng Liu)....Pages 593-605
Fabrication of Multiphase Particles and Grain Refinement of Al-Containing Magnesium (Chun-Hua Li, Yu Fu, Han Wang, Hai Hao)....Pages 607-614
Investigation on the Microstructure of Mg–Gd–Zn Alloy (Yu Liu, Dongqing Zhao, Hongtao Liu, Tao Lin, Yunteng Liu, Jixue Zhou)....Pages 615-621
The Evaluation Technology of Superalloy Cleanliness (Hua-Xia Zhang, Guo-Hong Ma, Tong-Jin Zhou, Wei Feng)....Pages 623-634
Effect of Microstructures on the Anisotropy of the Roll Casting Strip of Cold Rolled 3003 Aluminum Alloy (Canpei Ding, Gecheng Yuan, Haibin Guo, Zhiyong Long, Qian Yuan)....Pages 635-644
Modeling of Constitutive Equation and Microstructure Evolution of New Wrought Superalloy GH4066 (Yanju Wang, Chonglin Jia, Xingwu Li, Aixue Sha)....Pages 645-654
Effect of the Addition of Nanosized Y2O3 on the Mechanical Properties of WC-Bronze Composites (Linkai He, Youhong Sun, Chi Zhang, Jinhao Wu, Qingnan Meng)....Pages 655-663
High Temperature Ductility and Industrial Control Technology of Ni-Base Superalloy GH90 (Xin-li Wen, Qing-quan Zhang, Chao-lei Zhang, Bo Jiang, Ya-zheng Liu)....Pages 665-678
Effects of Mg/Si Ratio on the Microstructure of Aging Al–Mg–Si Alloys Containing Er and Zr Elements (Ning Li, Bolong Li, Tongbo Wang, Zuoren Nie)....Pages 679-687
Electrochemical Corrosion Properties of Ti46Zr20V12Cu5Be17 In Situ Metallic Glass Matrix Composites in HCl Solutions (Fan Yang)....Pages 689-695
Effect of Solution Treatment Temperature on the Mechanical Properties and Fracture Behavior of 7N01/7050 Aluminum Alloy Multilayer Plate (Guochuan Zhu, Shuhui Huang, Xiwu Li, Youzhi Tong, Zhihui Li, Baiqing Xiong et al.)....Pages 697-709
X-ray Investigations on Microstructural Evolution of Recrystallization of Single-Crystal Alloy DD6 (J. C. Xiong, J. R. Li, J. Yu)....Pages 711-717
Influence of Minor Grain Boundary Elements on the Solidification Behavior of a Re-containing Single-Crystal Superalloy (Jian Zhang, Jingxuan Zhao, Xiaotie Zhang, Yan Yang, Hao Chen, Hua Jiang et al.)....Pages 719-727
Back Matter ....Pages 729-733

Citation preview

Springer Proceedings in Physics 217

Yafang Han   Editor

Physics and Engineering of Metallic Materials Proceedings of Chinese Materials Conference 2018

Springer Proceedings in Physics Volume 217

The series Springer Proceedings in Physics, founded in 1984, is devoted to timely reports of state-of-the-art developments in physics and related sciences. Typically based on material presented at conferences, workshops and similar scientific meetings, volumes published in this series will constitute a comprehensive up-to-date source of reference on a field or subfield of relevance in contemporary physics. Proposals must include the following: – – – – –

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More information about this series at http://www.springer.com/series/361

Yafang Han Editor

Physics and Engineering of Metallic Materials Proceedings of Chinese Materials Conference 2018

123

Editor Yafang Han Chinese Materials Research Society Beijing, China

ISSN 0930-8989 ISSN 1867-4941 (electronic) Springer Proceedings in Physics ISBN 978-981-13-5943-9 ISBN 978-981-13-5944-6 (eBook) https://doi.org/10.1007/978-981-13-5944-6 Library of Congress Control Number: 2018966383 © Springer Nature Singapore Pte Ltd. 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 Singapore Pte Ltd. The registered company address is: 152 Beach Road, #21-01/04 Gateway East, Singapore 189721, Singapore

Preface

This is the proceedings of the selected papers presented at Chinese Materials Conference 2018 (CMC 2018) held in Xiamen City, Fujian, China, July 12–16, 2018. The Chinese Materials Conference (CMC) is the most important serial conference of Chinese Materials Research Society (C-MRS) and is held each year since early 1990s. Chinese Materials Conference 2018 had 35 Symposia covering four fields of Energy and environmental materials, Advanced functional materials, High-performance structural materials, and Design, preparation and characterization of materials. More than 5500 participants attended the conference, and the organizers received more than 500 technical papers. By recommendation of symposium organizers and after peer-reviewing, 94 papers are published in the present proceedings, which are divided into two volumes of Part 1: Physics and Engineering of Metallic Materials Part 2: Physics and Techniques of Ceramic and Polymeric Materials This is the volume for Part 1 including 70 papers selected from 7 symposia of Powder metallurgy, Advanced aluminum alloys, Advanced magnesium alloys, Superalloys, Metal matrix composites, Space materials science and technology, Nanoporous metal materials. The editor would like to give the thanks to the symposium chairs and all the paper reviewers of this volume, especially to Prof. Yafang Han, Prof. Xinqing Zhao, Prof. Jiangbo Sha, Prof. Chungen Zhou, Prof. Wenlong Xiao, and Prof. Tong Liu for their English polishing of the manuscripts. Beijing, China

Yafang Han

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Contents

Microstructures and Mechanical Properties of Extruded and Aged Mg–4Zn–2Al–2Sn–(0.6Mn) Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Dongqing Zhao, Yuansheng Yang, Xiaocun Song, Yu Liu, Cuicui Sun and Jixue Zhou Microstructural Evolution of an Al–Zn–Mg–Cu Aluminum Alloy During an Optimized Two-Step Homogenization Treatment . . . . . . . . . Hongwei Yan, Xiwu Li, Zhihui Li, Shuhui Huang, Hongwei Liu, Lizhen Yan, Wen Kai, Yong’an Zhang and Baiqing Xiong Microstructural Evolution and Phase Transformation of Al–Mg–Si Alloy Containing 3% Li During Homogenization . . . . . . . . . . . . . . . . . . Xiaokun Yang, Baiqing Xiong, Xiwu Li, Lizhen Yan, Zhihui Li, Yong’an Zhang, Hongwei Liu, Shuhui Huang, Hongwei Yan and Kai Wen Study on Stabilization Treatment of Al–Mg Alloy 5E83-H112 . . . . . . . . Xia Wu, Hui Huang, Shengping Wen, Xiaolan Wu, Zuoren Nie, Kunyuan Gao, Wei Wang and Bolong Li Simulation and Experimental Study on Hot Forging Process of SiCp/2A14 Composite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Hansen Zheng, Zhifeng Zhang and Yuelong Bai Effect of Boron Addition Methods on Microstructure and Mechanical Properties of a Near-a Titanium Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . Yingying Liu, Lihua Chai, Xiaozhao Ma, Yapeng Cui, Ziyong Chen and Zhilei Xiang Effect of Ultrasonic Treatment on Microstructure and Properties of Aluminum Alloy Rod Prepared by Continuous Casting and Rolling . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nan Zhou, Shuncheng Wang, Zhimin Zhang, Yuanhui Zhai and Guangyao Zhong

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Influence of High-Temperature Compression on Microstructure and Properties of Sintered Molybdenum . . . . . . . . . . . . . . . . . . . . . . . . Fu Wang, Zenglin Zhou, Yan Li, Zhilin Hui, Xueliang He and Xiaoying Fu Microstructure and Mechanical Properties of a Ti–Al–Sn–Zr–Mo–Nb–W–Si High-Temperature Titanium Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Y. W. Diao, X. Y. Song, W. J. Zhang, M. Y. Zhao, W. J. Ye and S. X. Hui Optimization of Cold-Rolling-Stabilization Process for High Mg-Containing Al Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Yueying Liang, Hui Huang, Xiaolan Wu, Shengping Wen, Kunyuan Gao, Zuoren Nie and Hongmei Li

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Flow Stress Behavior and Microstructural Evolution of a High-Alloying Al–Zn–Mg–Cu Alloy . . . . . . . . . . . . . . . . . . . . . . . . 105 Guohui Shi, Yong’an Zhang, Xiwu Li, Shuhui Huang, Zhihui Li, Lizhen Yan, Hongwei Yan and Hongwei Liu Influence of Space Environment on the Properties of Diamond/Cu Composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 117 Zhongnan Xie, Hong Guo and Ximin Zhang The Effects of Electron Beam Welding Parameters on Microstructure and Properties of GH4738 Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 123 Xinxu Li, Yong Zhang, Peihuan Li and Shaomin Lv Microstructure and Grain Refining Performance of High-Quality Al–5Ti–1B Master Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 133 Yuehua Kang, Shuncheng Wang, Nan Zhou and Dongfu Song Effect of Cu Content on Microstructure and Properties of Al–Mg–Si Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 143 Hong-Xiang Li, Shengli Guo, Peng Du and Sheng-Pu Liu Corrosion Behavior of Inconel 625 Alloy in Na2SO4–K2SO4 at High Temperature . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 153 Yuan-Jun Ma, Yutian Ding, Jian-Jun Liu, Yu-Bi Gao and Dong Zhang Effects of Extrusion Conditions on Microstructure and Age-Hardening Behaviors of Al–Zn–Mg Alloy . . . . . . . . . . . . . . . . . . . . 167 Y. L. Wang, H. C. Jiang, D. Zhang and L. J. Rong Process Optimization Design of High-Strength Ag–Cu–Ni Alloy Based on Orthogonal Experiments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 183 Helong Hu, Haibin Li, Wenjun Yu, Yongzhen Jiao, Tingyi Dong and Baoguo Lv

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Strain Distribution and Metal Flow of Bulk Forming of Molybdenum . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 191 Yihang Yang, Ailong Zheng, Zhimin Huang, Fusheng Peng and Houan Zhang Cold Deformation Behavior and Mechanical Properties of Forged Pure Nickel N6 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 201 Zexi Gao, Zhi Jia, Jinjin Ji, Dexue Liu and Yutian Ding Microstructure and Properties of Graphene/Copper Matrix Composites Prepared by In Situ Reduction . . . . . . . . . . . . . . . . . . . . . . 211 Xu Ran, Yutong Wang and Yong Wang The Effect of Aluminum Content on the Microstructures of Single-Phase c-TiAl-Based Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . 221 Yaodong Xuanyuan, Yan Long, Yinbiao Yan and Sen Yang Effect of Al–Er–Zr Master Alloy on Grain Refinement After Heat Treatment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 231 Haiyue Yu, Hui Huang, Zuoren Nie, Shengping Wen, Kunyuan Gao and Wei Wang Effect of In Situ TiB2 Particle Content on Microstructure and Properties of Cast Al–Si Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 241 Hongda Wang, Lihua Chai and He Li The Influence of Different Casting Methods on the Corrosion Resistance of Mg–Gd–Y Alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 249 Renju Cheng, Hong Tang, Wenjun Liu, Na Zhang, Hanwu Dong, Bin Jiang and Fusheng Pan Effect of Cu Additions and Extrusion Treatment on the Microstructure and Mechanical Properties of Mg–6Sn–1Al Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 257 Zhijian Ye, Tong Li, Gui Lou, Jianhang Yue and Xinying Teng Strain Rate Sensitivity of GH4720LI Alloy with Two Initial Microstructures During Hot Deformation . . . . . . . . . . . . . . . . . . . . . . . 273 Zhi-Peng Wan, Tao Wang, Yu Sun, Lian-Xi Hu, Zhao Li, Peihuan Li and Yong Zhang Microstructure and Properties of Rare-Earth B4C–Copper Composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 285 Yan Li, Meihui Song, Yu Zhang, Yanchun Li and Xiaochen Zhang Evaluation of Microstructure and Refining Effect of Al–TiB2 and Al–5Ti–1B Grain Refiners . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 295 Mengke He, Lihua Chai, Hongda Wang, Ziyong Chen and Yapeng Cui

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Effects of Graphene Content and Aging Process on Mechanical Properties and Corrosion Performance of an A356.2 Aluminum Matrix Composite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 307 Kang Wang, Jinfeng Leng, Ran Wang and Shaochen Zhang Microstructure and Performance of WAAM TiB2-Reinforced Al–Si-Based Composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 321 Yaqi Deng, Xianfeng Li, Liang Wu, Qingfeng Yang and Yanchi Chen Characterization of Graphene/Cu Composites Prepared by CVD and SPS . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 329 Xudong Wang, Haiping Zhang, Jiongli Li, Zhen Cao, Yue Wu and Na Li Hot Deformation Behavior of the As-Cast Mg–10.13Li–2.83Zn–2.78Al–0.13Si Alloy . . . . . . . . . . . . . . . . . . . . . . . . 337 Shouyang Gao, Defu Li, Peng Du and Shengli Guo Fatigue Behavior of Thermal Barrier Coated DD6 Single Crystal Superalloy at 900 °C . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 347 Jianmin Dong, Jiarong Li, Rende Mu and He Tian Microstructure and Room-Temperature Fracture Toughness of Nb–Ti–Si In Situ Composite Prepared by Selective Laser Melting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 357 Yongwang Kang, Fengwei Guo and Ming Li Simulation and Experimental Study on the Directional Solidification Process of a Single-Crystal Superalloy Plate Casting . . . . . . . . . . . . . . . 365 Runnan Wang, Qingyan Xu and Baicheng Liu Influence of Mo Content on the Microstructure Stability and Stress Rupture Properties of a Single Crystal Superalloy . . . . . . . . . . . . . . . . . 377 Z. X. Shi, S. Z. Liu, X. G. Wang and J. R. Li Preparation and Compressive Properties of Advanced Pore Morphology (APM) Foam Elements . . . . . . . . . . . . . . . . . . . . . . . . . . . . 383 Yanli Wang, Lucai Wang, Hong Xu and Qiaoyu Guo Effect of Temperature and Strain Amplitude on Low Cycle Fatigue Behaviour of DD11 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 391 Yuanyuan Guo, Yunsong Zhao, Jian Zhang, Yanfei Liu, Zhenyu Yang, Jiang Hua and Yushi Luo Oxidation Mechanism of Nb–Si-Based Ultra-High Temperature Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 399 Fengwei Guo, Yongwang Kang, Chenbo Xiao, Ming Li and Meiling Wu

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Microstructure and Mechanical Properties of Aluminum Alloy with Ultra-high Strength Prepared by Spray Forming . . . . . . . . . . . . . . 411 Shuhui Huang, Baiqing Xiong, Yong’an Zhang, Zhihui Li, Xiwu Li, Hongwei Liu, Hongwei Yan, Lizhen Yan and Kai Wen Effect of Pre-aging Technology on Microstructure and Mechanical Properties of 6111 Aluminum Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . 419 Hongwei Liu, Shuhui Huang, Baiqing Xiong, Yong’an Zhang, Zhihui Li, Xiwu Li, Hongwei Yan, Lizhen Yan and Kai Wen Effect of Zr Content on the Microstructure, Mechanical Properties, and Corrosion Resistance of Ti–27Nb–xZr Alloys . . . . . . . . . . . . . . . . . 429 Ying Xu, Huanhuan Wang, Yanqing Cai and Ziyan Wei Preparation of NiO by Precipitation Transformation and Its Supercapacitor Performance . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 439 Xinglei Wang, Yunqing Liu, Fan Zhang and Xiuling Yan Study on Fabrication and Compressive Properties of Mg/Al-Ordered Structure Composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 453 Han Wang, Yu Fu, Mingming Su and Hai Hao Study on Heat Treatment Process of Tungsten-Plated Diamond . . . . . . 463 Zhen Zhang, Hong Guo, Zhongnan Xie and Ximin Zhang Effect of Sc Modification and Pulping Process on Semi-Solid Structure of A356 Aluminum Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 471 Yuxin Zhang, Hengbin Liao, Yong Dong, Anfu Chen, Xiaoling Fu and Zhengrong Zhang Densification Behavior of High-Volume-Fraction (B4C + W)Reinforced Aluminum Matrix Composites Prepared by Hot Pressing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 485 Tiantian Guo, Changhui Mao, Shuwang Ma, Zheng Lv and Guosong Zhang The Finite Element Simulation of Diffusion Bonding for TiAl/Ti2AlNb Annular Structural Component . . . . . . . . . . . . . . . . . 495 Xiaoqiang Zhang, Bin Tang, Jinshan Li and Hongchao Kou Effect of Hf Content on the Stress Rupture Properties of the Second-Generation Single-Crystal Superalloy with LAB . . . . . . . . . . . . 505 Yipeng Zhang, Jianchao Qin, Zhaohui Huang, Renjie Cui and Kai Guan High-Cycle Fatigue of Mg–6Gd–3Y–0.5Zr Cast Magnesium Alloys . . . . 515 Dehao Meng, Peijie Li, Duanzhi Wang, Wenquan Yuan and Wencai Liu Superplastic Forming Properties and Instability of Magnesium Alloy Sheet . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 527 Meijuan Song, Chuanhui Huang, Wenchao Jiang and Weina Lu

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The Influence of Fused Mullite Mineralizer on Mechanical Properties of Ceramic Shell . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 537 Longpei Dong, Jiansheng Yao, Bin Shen, Feipeng Guo, Xiaoyu Li and Dingzhong Tang Influences of Welding Current on Joint Structure and Mechanical Property of AZ31B Magnesium Board . . . . . . . . . . . . . . . . . . . . . . . . . . 547 Yingchun Tian, Ping Gao, Jianjun Lv and Mingyi Zhang Surface Characterization of Nickel-Base Superalloy Powder . . . . . . . . . 561 Wenyong Xu, Yufeng Liu, Hua Yuan, Zhou Li and Guoqing Zhang Microstructure and Mechanical Properties of 6005A-T5 Aluminum Alloy Welded Joints by Friction Stir Welding and Metal Inert Gas Welding . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 569 Jingxuan Liu, Jian Shen, Xiwu Li, Lizhen Yan, Hongwei Yan, Hongwei Liu, Zhihui Li, Yong’an Zhang and Baiqing Xiong Deformation Behavior and Constitutive Model for Isothermal Compression of TC4 Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 581 Jiangwei Zhong, Pan Tao, Qingyan Xu, Baicheng Liu and Zhijun Ji Recent Development of Hot-Pressed-/Deformed Nd–Fe–B Permanent Magnets . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 593 Ke Lv, Rui Dong, Mingjing Zhao and Guozheng Liu Fabrication of Multiphase Particles and Grain Refinement of Al-Containing Magnesium . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 607 Chun-Hua Li, Yu Fu, Han Wang and Hai Hao Investigation on the Microstructure of Mg–Gd–Zn Alloy . . . . . . . . . . . . 615 Yu Liu, Dongqing Zhao, Hongtao Liu, Tao Lin, Yunteng Liu and Jixue Zhou The Evaluation Technology of Superalloy Cleanliness . . . . . . . . . . . . . . 623 Hua-Xia Zhang, Guo-Hong Ma, Tong-Jin Zhou and Wei Feng Effect of Microstructures on the Anisotropy of the Roll Casting Strip of Cold Rolled 3003 Aluminum Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . 635 Canpei Ding, Gecheng Yuan, Haibin Guo, Zhiyong Long and Qian Yuan Modeling of Constitutive Equation and Microstructure Evolution of New Wrought Superalloy GH4066 . . . . . . . . . . . . . . . . . . . . . . . . . . . 645 Yanju Wang, Chonglin Jia, Xingwu Li and Aixue Sha Effect of the Addition of Nanosized Y2O3 on the Mechanical Properties of WC-Bronze Composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 655 Linkai He, Youhong Sun, Chi Zhang, Jinhao Wu and Qingnan Meng

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High Temperature Ductility and Industrial Control Technology of Ni-Base Superalloy GH90 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 665 Xin-li Wen, Qing-quan Zhang, Chao-lei Zhang, Bo Jiang and Ya-zheng Liu Effects of Mg/Si Ratio on the Microstructure of Aging Al–Mg–Si Alloys Containing Er and Zr Elements . . . . . . . . . . . . . . . . . . . . . . . . . 679 Ning Li, Bolong Li, Tongbo Wang and Zuoren Nie Electrochemical Corrosion Properties of Ti46Zr20V12Cu5Be17 In Situ Metallic Glass Matrix Composites in HCl Solutions . . . . . . . . . . . . . . . . 689 Fan Yang Effect of Solution Treatment Temperature on the Mechanical Properties and Fracture Behavior of 7N01/7050 Aluminum Alloy Multilayer Plate . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 697 Guochuan Zhu, Shuhui Huang, Xiwu Li, Youzhi Tong, Zhihui Li, Baiqing Xiong and Yong’an Zhang X-ray Investigations on Microstructural Evolution of Recrystallization of Single-Crystal Alloy DD6 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 711 J. C. Xiong, J. R. Li and J. Yu Influence of Minor Grain Boundary Elements on the Solidification Behavior of a Re-containing Single-Crystal Superalloy . . . . . . . . . . . . . 719 Jian Zhang, Jingxuan Zhao, Xiaotie Zhang, Yan Yang, Hao Chen, Hua Jiang and Li She Index . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 729

Contributors

Yuelong Bai General Research Institute for Nonferrous Metals, Beijing, China Yanqing Cai College of Material Science and Engineering, North China University of Science and Technology, Tangshan, China Zhen Cao AECC Beijing Institute of Aeronautical Materials, Beijing, China Beijing Institute of Graphene Technology, Beijing, China Beijing Engineering Research Centre of Graphene Application, Beijing, China Lihua Chai College of Materials Science and Engineering, Beijing University of Technology, Beijing, China Anfu Chen School of Materials and Energy, Guangdong University of Technology, Guangzhou, Guangdong, China Hao Chen Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Yanchi Chen Shanghai JiaoTong University, Shanghai, China Ziyong Chen College of Materials Science and Engineering, Beijing University of Technology, Beijing, China Renju Cheng Chongqing Academy of Science and Technology, Chongqing, China Renjie Cui Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Yapeng Cui College of Materials Science and Engineering, Beijing University of Technology, Beijing, China Yaqi Deng Shanghai JiaoTong University, Shanghai, China Y. W. Diao State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co. Ltd., Beijing, China

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xvi

Contributors

Canpei Ding School of Materials and Energy, Guangdong University of Technology, Guangzhou, China Yutian Ding State Key Laboratory of Advanced Processing and Recycling of Nonferrous Metals, Lanzhou University of Technology, Lanzhou, Gansu Province, China Hanwu Dong Chongqing Academy of Science and Technology, Chongqing, China Jianmin Dong Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Longpei Dong Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Rui Dong University of Science and Technology Inner Mongolia, Baotou, China Tingyi Dong Beijing Trillion Metals Co., Ltd., Beijing, China GRIKIN Advanced Material Co., Ltd., Beijing, China Yong Dong School of Materials and Energy, Guangdong University of Technology, Guangzhou, Guangdong, China Peng Du General Research Institute for Non-ferrous Metals, Beijing, China Wei Feng AECC Beijing Institute of Aeronautical Materials, Beijing, China Xiaoling Fu School of Materials and Energy, Guangdong University of Technology, Guangzhou, Guangdong, China Xiaoying Fu Powder Metallurgy and Special Materials Research Department, GRIMAT Engineering Institute Co., Ltd, Beijing, China Yu Fu School of Materials Science and Engineering, Dalian University of Technology, Dalian, China Kunyuan Gao School of Materials Science and Engineering, Beijing University of Technology, Beijing, People’s Republic of China Ping Gao Inner Mongolia Institute of Metal Materials, Baotou, Inner Mongolia, China Shouyang Gao General Research Institute for Nonferrous Metals, Beijing, China Yu-Bi Gao State Key Laboratory of Advanced Processing and Recycling of Non-ferrous Metals, Lanzhou University of Technology, Lanzhou, China Zexi Gao School of Material Science and Engineering, Lanzhou University of Technology, Lanzhou, Gansu, China

Contributors

xvii

Kai Guan Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Feipeng Guo Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Fengwei Guo Science & Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Haibin Guo School of Materials and Energy, Guangdong University of Technology, Guangzhou, China Hong Guo National Engineering Research Center for Nonferrous Metals Composites, General Research Institute for Non-ferrous Metals, Beijing, China Qiaoyu Guo Taiyuan University of Science and Technology, Taiyuan, Shanxi Province, China Shengli Guo General Research Institute for Nonferrous Metals, Beijing, China Tiantian Guo Advanced Electronic Materials Institute, GRIMAT Engineering Institute Co., Ltd., Beijing, China Yuanyuan Guo Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Hai Hao School of Materials Science and Engineering, Dalian University of Technology, Dalian, China Linkai He College of Construction Engineering, Jilin University, Changchun, People’s Republic of China Key Lab of Drilling and Exploitation Technology in Complex Conditions, Ministry of Land and Resources, Changchun, People’s Republic of China Mengke He College of Materials Science and Engineering, Beijing University of Technology, Beijing, China Xueliang He Powder Metallurgy and Special Materials Research Department, GRIMAT Engineering Institute Co., Ltd, Beijing, China Helong Hu Beijing Trillion Metals Co., Ltd., Beijing, China Lian-Xi Hu National Key Laboratory for Precision Hot Processing of Metals, Harbin Institute of Technology, Harbin, China Jiang Hua Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Chuanhui Huang School of Mechanical and Electrical Engineering, Xuzhou Institute of Technology, Xuzhou, People’s Republic of China

xviii

Contributors

Hui Huang School of Materials Science and Engineering, Beijing University of Technology, Beijing, People’s Republic of China Shuhui Huang State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co., Ltd., Beijing, China State Key Laboratory of Nonferrous Metals and Processes, GRINM Group Co., Ltd., Beijing, China State Key Laboratory of Nonferrous Metals and Processes, General Research Institute for Nonferrous Metals, Beijing, China Zhaohui Huang Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Zhimin Huang Department of R&D, Xiamen Honglu Tungsten Molybdenum Industry Co. Ltd., Xiamen, China S. X. Hui State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co. Ltd., Beijing, China Zhilin Hui Powder Metallurgy and Special Materials Research Department, GRIMAT Engineering Institute Co., Ltd, Beijing, China Jinjin Ji School of Materials Engineering, Lanzhou Institute of Technology, Lanzhou, Gansu, China Zhijun Ji AECC Beijing Institute of Aeronautical Materials, Beijing, China Chonglin Jia Aviation Engine Corporation of China, Beijing Institute of Aeronautical Materials, Beijing, China Zhi Jia State Key Laboratory of Advanced Processing and Recycling of Nonferrous Metals, Lanzhou University of Technology, Lanzhou, Gansu Province, China School of Material Science and Engineering, Lanzhou University of Technology, Lanzhou, Gansu, China Bin Jiang College of Materials Science and Engineering, Chongqing University, Chongqing, China National Engineering Research Center for Magnesium Alloy, Chongqing University, Chongqing, China Bo Jiang School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing, China H. C. Jiang CAS Key Laboratory of Nuclear Materials and Safety Assessment, Institute of Metal Research, Chinese Academy of Science, Shenyang, China Hua Jiang Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China

Contributors

xix

Wenchao Jiang School of Mechanical and Electrical Engineering, Xuzhou Institute of Technology, Xuzhou, People’s Republic of China Yongzhen Jiao Beijing Trillion Metals Co., Ltd., Beijing, China Wen Kai State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co., Ltd., Beijing, China Yongwang Kang Science & Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Yuehua Kang Guangdong Institute of Materials and Processing, Guangzhou, China Hongchao Kou State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an, China Jinfeng Leng School of Materials Science and Engineering, University of Jinan, Jinan, People’s Republic of China Bolong Li School of Materials Science and Engineering, Beijing University of Technology, Beijing, People’s Republic of China Chun-Hua Li Dalian University of Technology, Dalian, China Defu Li General Research Institute for Nonferrous Metals, Beijing, China Haibin Li Beijing Trillion Metals Co., Ltd., Beijing, China GRIKIN Advanced Material Co., Ltd., Beijing, China He Li College of Materials Science and Engineering, Beijing University of Technology, Beijing, China Hong-Xiang Li General Research Institute for Non-ferrous Metals, Beijing, China Hongmei Li School of Materials Science and Engineering, Beijing University of Technology, Beijing, People’s Republic of China J. R. Li Science and Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Jiarong Li Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Jinshan Li State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an, China Jiongli Li AECC Beijing Institute of Aeronautical Materials, Beijing, China Beijing Institute of Graphene Technology, Beijing, China Beijing Engineering Research Centre of Graphene Application, Beijing, China

xx

Contributors

Ming Li Science & Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Na Li AECC Beijing Institute of Aeronautical Materials, Beijing, China Ning Li Beijing University of Technology, Chaoyang District, Beijing, China Peihuan Li Science and Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Peijie Li Tsinghua University, Beijing, China Tong Li School of Materials Science and Engineering, University of Jinan, Jinan, People’s Republic of China Xianfeng Li Shanghai JiaoTong University, Shanghai, China Xiaoyu Li Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Xingwu Li Aviation Engine Corporation of China, Beijing Institute of Aeronautical Materials, Beijing, China Xinxu Li Science and Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China School of Metallurgy, Northeastern University, Shenyang, China Xiwu Li State Key Laboratory of Nonferrous Metals and Processes, GRINM Group Co., Ltd., Beijing, China State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co., Ltd., Beijing, China State Key Laboratory of Nonferrous Metals and Processes, General Research Institute for Nonferrous Metals, Beijing, China Yan Li Powder Metallurgy and Special Materials Research Department, GRIMAT Engineering Institute Co., Ltd, Beijing, China Institute of Advanced Technology, Heilongjiang Academy of Sciences, Harbin, China Yanchun Li Institute of Advanced Technology, Heilongjiang Academy of Sciences, Harbin, China Zhao Li Science and Technology on Advanced High Temperature Structural Materials Laboratory, AEEC Beijing Institute of Aeronautical Materials, Beijing, China Zhihui Li State Key Laboratory of Nonferrous Metals and Processes, General Research Institute for Nonferrous Metals, GRINM Group Co., Ltd., Beijing, China

Contributors

xxi

Zhou Li Science and Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Yueying Liang School of Materials Science and Engineering, Beijing University of Technology, Beijing, People’s Republic of China Hengbin Liao School of Materials and Energy, Guangdong University of Technology, Guangzhou, Guangdong, China Tao Lin Advanced Materials Institute, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Shandong Provincial Key Laboratory for High Strength Lightweight Metallic Materials, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Baicheng Liu Key Laboratory for Advanced Materials Processing Technology (MOE), School of Materials Science and Engineering, Tsinghua University, Beijing, China Dexue Liu State Key Laboratory of Advanced Processing and Recycling of Nonferrous Metals, Lanzhou University of Technology, Lanzhou, Gansu Province, China Guozheng Liu Baotou Research Institute of Rare Earth, Baotou, China Hongtao Liu Advanced Materials Institute, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Shandong Provincial Key Laboratory for High Strength Lightweight Metallic Materials, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Hongwei Liu State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co., Ltd., Beijing, China Jian-Jun Liu State Key Laboratory of Advanced Processing and Recycling of Non-ferrous Metals, Lanzhou University of Technology, Lanzhou, China Jingxuan Liu State Key Laboratory of Nonferrous Metals and Processes, GRINM Group Co., Ltd, Beijing, China S. Z. Liu Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Sheng-Pu Liu General Research Institute for Non-ferrous Metals, Beijing, China Wencai Liu National Engineering Research Center of Light Alloy Net Forming and State Key Laboratory of Metal Matrix Composites, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai, China Wenjun Liu Chongqing Academy of Science and Technology, Chongqing, China

xxii

Contributors

Ya-zheng Liu School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing, China Yanfei Liu Sichuan Gas Turbine Research Establishment, Sichuan, China Yingying Liu College of Materials Science and Engineering, Beijing University of Technology, Beijing, China Yu Liu Advanced Materials Institute, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Shandong Provincial Key Laboratory for High Strength Lightweight Metallic Materials, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Yufeng Liu Science and Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Yunqing Liu Laboratory of Condensed Matter Phase Transition and Microstructure, School of Chemistry and Environmental Sciences, Yili Normal University, Xinjiang, Yining, China Yunteng Liu Advanced Materials Institute, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Shandong Provincial Key Laboratory for High Strength Lightweight Metallic Materials, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Yan Long School of Material Science and Engineering, Nanjing University of Science and Technology, Nanjing, China Zhiyong Long School of Materials and Energy, Guangdong University of Technology, Guangzhou, China Gui Lou School of Materials Science and Engineering, University of Jinan, Jinan, People’s Republic of China Weina Lu School of Mechanical and Electrical Engineering, Xuzhou Institute of Technology, Xuzhou, People’s Republic of China Yushi Luo Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Baoguo Lv Beijing Trillion Metals Co., Ltd., Beijing, China GRIKIN Advanced Material Co., Ltd., Beijing, China Jianjun Lv Inner Mongolia Institute of Metal Materials, Baotou, Inner Mongolia, China Ke Lv Baotou Research Institute of Rare Earth, Baotou, China

Contributors

xxiii

Shaomin Lv Science and Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Institute for Advanced Materials and Technology, University of Science and Technology Beijing, Beijing, China Zheng Lv Advanced Electronic Materials Institute, GRIMAT Engineering Institute Co., Ltd., Beijing, China Guo-Hong Ma AECC Beijing Institute of Aeronautical Materials, Beijing, China Shuwang Ma Advanced Electronic Materials Institute, GRIMAT Engineering Institute Co., Ltd., Beijing, China Xiaozhao Ma College of Materials Science and Engineering, Beijing University of Technology, Beijing, China Yuan-Jun Ma State Key Laboratory of Advanced Processing and Recycling of Non-ferrous Metals, Lanzhou University of Technology, Lanzhou, China Changhui Mao Advanced Electronic Materials Institute, GRIMAT Engineering Institute Co., Ltd., Beijing, China Dehao Meng Tsinghua University, Beijing, China Beijing Institute of Aerospace Systems Engineering, Beijing, China Qingnan Meng College of Construction Engineering, Jilin University, Changchun, People’s Republic of China Key Lab of Drilling and Exploitation Technology in Complex Conditions, Ministry of Land and Resources, Changchun, People’s Republic of China Rende Mu Aviation Key Laboratory of Science and Technology on Advanced Corrosion and Protection for Aviation Material, Beijing Institute of Aeronautical Materials, Beijing, China Zuoren Nie School of Materials Science and Engineering, Beijing University of Technology, Beijing, People’s Republic of China Beijing University of Technology, Chaoyang District, Beijing, China Fusheng Pan Chongqing Academy of Science and Technology, Chongqing, China College of Materials Science and Engineering, Chongqing University, Chongqing, China National Engineering Research Center for Magnesium Alloy, Chongqing University, Chongqing, China Fusheng Peng Department of R&D, Xiamen Honglu Tungsten Molybdenum Industry Co. Ltd., Xiamen, China Jianchao Qin Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China

xxiv

Contributors

Xu Ran Key Laboratory of Advanced Structural Materials, Ministry of Education, Changchun University of Technology, Changchun, Jilin, China L. J. Rong CAS Key Laboratory of Nuclear Materials and Safety Assessment, Institute of Metal Research, Chinese Academy of Science, Shenyang, China Aixue Sha Aviation Engine Corporation of China, Beijing Institute of Aeronautical Materials, Beijing, China Li She Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Bin Shen Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Jian Shen State Key Laboratory of Nonferrous Metals and Processes, GRINM Group Co., Ltd, Beijing, China Guohui Shi State Key Laboratory of Nonferrous Metals and Processes, GRINM Group Co., Ltd., Beijing, China Z. X. Shi Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Dongfu Song Guangdong Institute of Materials and Processing, Guangzhou, China Meihui Song Institute of Advanced Technology, Heilongjiang Academy of Sciences, Harbin, China Meijuan Song School of Mechanical and Electrical Engineering, Xuzhou Institute of Technology, Xuzhou, People’s Republic of China X. Y. Song State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co. Ltd., Beijing, China Xiaocun Song Shandong Provincial Key Laboratory for High Strength Lightweight Metallic Materials, Advanced Materials Institute, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Shandong Engineering Research Center for Lightweight Automobiles Magnesium Alloys, Advanced Materials Institute, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Mingming Su School of Materials Science and Engineering, Dalian University of Technology, Dalian, China Cuicui Sun Shandong Provincial Key Laboratory for High Strength Lightweight Metallic Materials, Advanced Materials Institute, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China

Contributors

xxv

Youhong Sun College of Construction Engineering, Jilin University, Changchun, People’s Republic of China Key Lab of Drilling and Exploitation Technology in Complex Conditions, Ministry of Land and Resources, Changchun, People’s Republic of China Yu Sun National Key Laboratory for Precision Hot Processing of Metals, Harbin Institute of Technology, Harbin, China Bin Tang State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an, China Dingzhong Tang Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Hong Tang College of Materials Science and Engineering, Chongqing University, Chongqing, China Pan Tao Key Laboratory for Advanced Materials Processing Technology, School of Materials Science and Engineering, Tsinghua University, Beijing, China Xinying Teng School of Materials Science and Engineering, University of Jinan, Jinan, People’s Republic of China He Tian Aviation Key Laboratory of Science and Technology on Advanced Corrosion and Protection for Aviation Material, Beijing Institute of Aeronautical Materials, Beijing, China Yingchun Tian Inner Mongolia Institute of Metal Materials, Baotou, Inner Mongolia, China Youzhi Tong Northeast Light Alloy Co., Ltd., Harbin, Heilongjiang, China Zhi-Peng Wan Science and Technology on Advanced High Temperature Structural Materials Laboratory, AEEC Beijing Institute of Aeronautical Materials, Beijing, China National Key Laboratory for Precision Hot Processing of Metals, Harbin Institute of Technology, Harbin, China Duanzhi Wang Beijing Institute of Aerospace Systems Engineering, Beijing, China Fu Wang Powder Metallurgy and Special Materials Research Department, GRIMAT Engineering Institute Co., Ltd, Beijing, China Han Wang School of Materials Science and Engineering, Dalian University of Technology, Dalian, China Hongda Wang College of Materials Science and Engineering, Beijing University of Technology, Beijing, China

xxvi

Contributors

Huanhuan Wang College of Material Science and Engineering, North China University of Science and Technology, Tangshan, China Kang Wang School of Materials Science and Engineering, University of Jinan, Jinan, People’s Republic of China Lucai Wang Taiyuan University of Science and Technology, Taiyuan, Shanxi Province, China Ran Wang School of Materials Science and Engineering, University of Jinan, Jinan, People’s Republic of China Runnan Wang Key Laboratory for Advanced Materials Processing Technology (MOE), School of Materials Science and Engineering, Tsinghua University, Beijing, China Shuncheng Wang Guangdong Institute of Materials and Processing, Guangzhou, China Tao Wang Science and Technology on Advanced High Temperature Structural Materials Laboratory, AEEC Beijing Institute of Aeronautical Materials, Beijing, China Tongbo Wang Beijing University of Technology, Chaoyang District, Beijing, China Wei Wang School of Materials Science and Engineering, Beijing University of Technology, Beijing, People’s Republic of China X. G. Wang Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Xinglei Wang Laboratory of Condensed Matter Phase Transition and Microstructure, School of Chemistry and Environmental Sciences, Yili Normal University, Xinjiang, Yining, China Xudong Wang AECC Beijing Institute of Aeronautical Materials, Beijing, China Beijing Institute of Graphene Technology, Beijing, China Beijing Engineering Research Centre of Graphene Application, Beijing, China Y. L. Wang CAS Key Laboratory of Nuclear Materials and Safety Assessment, Institute of Metal Research, Chinese Academy of Science, Shenyang, China School of Materials Science and Engineering, University of Science and Technology of China, Hefei, China Yanju Wang Aviation Engine Corporation of China, Beijing Institute of Aeronautical Materials, Beijing, China Yanli Wang North University of China, Taiyuan, Shanxi Province, China Yong Wang Key Laboratory of Advanced Structural Materials, Ministry of Education, Changchun University of Technology, Changchun, Jilin, China

Contributors

xxvii

Yutong Wang Key Laboratory of Advanced Structural Materials, Ministry of Education, Changchun University of Technology, Changchun, Jilin, China Ziyan Wei College of Material Science and Engineering, North China University of Science and Technology, Tangshan, China Kai Wen State Key Laboratory of Nonferrous Metals and Processes, GRINM Group Co. Ltd., Beijing, China State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co., Ltd., Beijing, China Shengping Wen School of Materials Science and Engineering, Beijing University of Technology, Beijing, People’s Republic of China Xin-li Wen Beijing Beiye Functional Materials Corporation, Beijing, China Jinhao Wu College of Construction Engineering, Jilin University, Changchun, People’s Republic of China Key Lab of Drilling and Exploitation Technology in Complex Conditions, Ministry of Land and Resources, Changchun, People’s Republic of China Liang Wu Shanghai JiaoTong University, Shanghai, China Meiling Wu Science & Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Xia Wu School of Materials Science and Engineering, Beijing University of Technology, Beijing, People’s Republic of China Xiaolan Wu School of Materials Science and Engineering, Beijing University of Technology, Beijing, People’s Republic of China Yue Wu AECC Beijing Institute of Aeronautical Materials, Beijing, China Beijing Institute of Graphene Technology, Beijing, China Beijing Engineering Research Centre of Graphene Application, Beijing, China Zhilei Xiang College of Materials Science and Engineering, Beijing University of Technology, Beijing, China Chenbo Xiao Science & Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Zhongnan Xie General Research Institute for Nonferrous Metals, Beijing, China National Engineering Research Center for Nonferrous Metals Composites, General Research Institute for Non-ferrous Metals, Beijing, China Baiqing Xiong State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co., Ltd., Beijing, China State Key Laboratory of Nonferrous Metals and Processes, GRINM Group Co., Ltd., Beijing, China

xxviii

Contributors

State Key Laboratory of Nonferrous Metals and Processes, General Research Institute for Nonferrous Metals, Beijing, China J. C. Xiong Science and Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Hong Xu North University of China, Taiyuan, Shanxi Province, China Qingyan Xu Key Laboratory for Advanced Materials Processing Technology, School of Materials Science and Engineering, Tsinghua University, Beijing, China Qingyan Xu Key Laboratory for Advanced Materials Processing Technology (MOE), School of Materials Science and Engineering, Tsinghua University, Beijing, China Wenyong Xu Science and Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Ying Xu College of Material Science and Engineering, North China University of Science and Technology, Tangshan, China Yaodong Xuanyuan School of Material Science and Engineering, Nanjing University of Science and Technology, Nanjing, China Hongwei Yan State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co., Ltd., Beijing, China State Key Laboratory of Nonferrous Metals and Processes, GRINM Group Co., Ltd., Beijing, China Lizhen Yan State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co., Ltd., Beijing, China State Key Laboratory of Nonferrous Metals and Processes, GRINM Group Co., Ltd., Beijing, China Xiuling Yan Laboratory of Condensed Matter Phase Transition and Microstructure, School of Chemistry and Environmental Sciences, Yili Normal University, Xinjiang, Yining, China Yinbiao Yan School of Material Science and Engineering, Nanjing University of Science and Technology, Nanjing, China Fan Yang Beijing Beiye Functional Materials Corporation, Haidian District, Beijing, China Qingfeng Yang Shanghai JiaoTong University, Shanghai, China Sen Yang School of Material Science and Engineering, Nanjing University of Science and Technology, Nanjing, China

Contributors

xxix

Xiaokun Yang State Key Laboratory of Nonferrous Metals and Processes, GRINM Group Co. Ltd., Beijing, China Yan Yang Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Yihang Yang Fujian Key Laboratory of Functional Materials and Applications, School of Materials Science and Engineering, Xiamen University of Technology, Xiamen, China Department of R&D, Xiamen Honglu Tungsten Molybdenum Industry Co. Ltd., Xiamen, China Yuansheng Yang Shandong Provincial Key Laboratory for High Strength Lightweight Metallic Materials, Advanced Materials Institute, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Institute of Metal Research, Chinese Academy of Sciences, Shenyang, China Zhenyu Yang Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Jiansheng Yao Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China W. J. Ye State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co. Ltd., Beijing, China Zhijian Ye School of Materials Science and Engineering, University of Jinan, Jinan, People’s Republic of China Haiyue Yu School of Materials Science and Engineering, Beijing University of Technology, Beijing, People’s Republic of China J. Yu Science and Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Wenjun Yu Beijing Trillion Metals Co., Ltd., Beijing, China Gecheng Yuan School of Materials and Energy, Guangdong University of Technology, Guangzhou, China Hua Yuan Science and Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Qian Yuan School of Materials and Energy, Guangdong University of Technology, Guangzhou, China Wenquan Yuan Beijing Institute of Aerospace Systems Engineering, Beijing, China

xxx

Contributors

Jianhang Yue School of Materials Science and Engineering, University of Jinan, Jinan, People’s Republic of China Yuanhui Zhai Guangdong Shine Cables Co., Ltd., Qingyuan, China Chao-lei Zhang School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing, China Chi Zhang College of Construction Engineering, Jilin University, Changchun, People’s Republic of China Key Lab of Drilling and Exploitation Technology in Complex Conditions, Ministry of Land and Resources, Changchun, People’s Republic of China D. Zhang CAS Key Laboratory of Nuclear Materials and Safety Assessment, Institute of Metal Research, Chinese Academy of Science, Shenyang, China School of Materials Science and Engineering, University of Science and Technology of China, Hefei, China Dong Zhang State Key Laboratory of Nickel and Cobalt Resources Comprehensive Utilization, Jinchuan Group Co., Ltd., Jinchang, China Fan Zhang Laboratory of Condensed Matter Phase Transition and Microstructure, School of Chemistry and Environmental Sciences, Yili Normal University, Xinjiang, Yining, China Guoqing Zhang Science and Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Guosong Zhang Advanced Electronic Materials Institute, GRIMAT Engineering Institute Co., Ltd., Beijing, China Haiping Zhang AECC Beijing Institute of Aeronautical Materials, Beijing, China Beijing Institute of Graphene Technology, Beijing, China Beijing Engineering Research Centre of Graphene Application, Beijing, China Houan Zhang Fujian Key Laboratory of Functional Materials and Applications, School of Materials Science and Engineering, Xiamen University of Technology, Xiamen, China Fujian Collaborative Innovation Center for R&D of Coach and Special Vehicle, Xiamen University of Technology, Xiamen, China Hua-Xia Zhang AECC Beijing Institute of Aeronautical Materials, Beijing, China Jian Zhang Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Mingyi Zhang Inner Mongolia Institute of Metal Materials, Baotou, Inner Mongolia, China

Contributors

xxxi

Na Zhang Chongqing Academy of Science and Technology, Chongqing, China Qing-quan Zhang Beijing Beiye Functional Materials Corporation, Beijing, China Shaochen Zhang School of Materials Science and Engineering, University of Jinan, Jinan, People’s Republic of China W. J. Zhang State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co. Ltd., Beijing, China Xiaochen Zhang Institute of Advanced Technology, Heilongjiang Academy of Sciences, Harbin, China Xiaoqiang Zhang State Key Laboratory of Solidification Northwestern Polytechnical University, Xi’an, China

Processing,

Xiaotie Zhang Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Ximin Zhang General Research Institute for Nonferrous Metals, Beijing, China National Engineering Research Center for Nonferrous Metals Composites, General Research Institute for Non-ferrous Metals, Beijing, China Yipeng Zhang Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Yong Zhang Science and Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing, China Yong’an Zhang State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co., Ltd., Beijing, China State Key Laboratory of Nonferrous Metals and Processes, GRINM Group Co. Ltd., Beijing, China Yu Zhang Institute of Advanced Technology, Heilongjiang Academy of Sciences, Harbin, China Yuxin Zhang School of Materials and Energy, Guangdong University of Technology, Guangzhou, Guangdong, China Zhen Zhang National Engineering Research Center for Nonferrous Metals Composites, General Research Institute for Non-ferrous Metals, Beijing, China Zhengrong Zhang School of Materials and Energy, Guangdong University of Technology, Guangzhou, Guangdong, China

xxxii

Contributors

Zhifeng Zhang General Research Institute for Nonferrous Metals, Beijing, China Zhimin Zhang Guangdong Shine Cables Co., Ltd., Qingyuan, China Dongqing Zhao Advanced Materials Institute, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Shandong Provincial Key Laboratory for High Strength Lightweight Metallic Materials, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Institute of Metal Research, Chinese Academy of Sciences, Shenyang, China Jingxuan Zhao Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China M. Y. Zhao State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co. Ltd., Beijing, China Mingjing Zhao Baotou Research Institute of Rare Earth, Baotou, China Yunsong Zhao Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing, China Ailong Zheng Department of R&D, Xiamen Honglu Tungsten Molybdenum Industry Co. Ltd., Xiamen, China Hansen Zheng General Research Institute for Nonferrous Metals, Beijing, China Guangyao Zhong Guangdong Shine Cables Co., Ltd., Qingyuan, China Jiangwei Zhong Key Laboratory for Advanced Materials Processing Technology, School of Materials Science and Engineering, Tsinghua University, Beijing, China Jixue Zhou Advanced Materials Institute, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Shandong Provincial Key Laboratory for High Strength Lightweight Metallic Materials, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Shandong Engineering Research Center for Lightweight Automobiles Magnesium Alloys, Advanced Materials Institute, Qilu University of Technology (Shandong Academy of Sciences), Jinan, China Guangdong Institute of Materials and Processing, Guangzhou, China Nan Zhou Guangdong Institute of Materials and Processing, Guangzhou, China Tong-Jin Zhou AECC Beijing Institute of Aeronautical Materials, Beijing, China

Contributors

xxxiii

Zenglin Zhou Powder Metallurgy and Special Materials Research Department, GRIMAT Engineering Institute Co., Ltd, Beijing, China Guochuan Zhu State Key Laboratory of Nonferrous Metals and Processes, General Research Institute for Nonferrous Metals, Beijing, China

Microstructures and Mechanical Properties of Extruded and Aged Mg–4Zn–2Al–2Sn–(0.6Mn) Alloy Dongqing Zhao, Yuansheng Yang, Xiaocun Song, Yu Liu, Cuicui Sun and Jixue Zhou

Abstract The previous study has shown that the mechanical properties of extruded Mg–Zn–Al–Sn alloys decrease with the increasing of extrusion temperature because of the coarsened microstructure. In order to improve the microstructure of Mg–Zn–Al–Sn alloy extruded at a high temperature, a small amount of Mn (0.6 wt%) was added to Mg–4Zn–2Al–2Sn (ZAT422, wt%) alloy in the present study, and ZAT422 and Mg–4Zn–2Al–2Sn–0.6Mn (ZATM4220, wt%) alloys were extruded into bars with 16 mm in diameter at an extrusion temperature of 648 K. Then the microstructures and mechanical properties of ZAT422 and ZATM4220 alloys were examined. It was found that the extruded ZATM4220 alloy exhibited finer deformation microstructure and better mechanical properties than ZAT422 alloy.

1 Introduction Compared with cast magnesium alloy, wrought magnesium alloy shows higher mechanical strength and better ductility due to their more refined microstructure resulted from the thermomechanical treatment. Therefore, there are some applications where wrought magnesium may be more qualified, such as some bearing structural parts on automotive vehicle and spacecraft, bicycle frame, shells of notebook computers, etc. [1]. However, the most frequently used commercial wrought magnesium alloys are no longer meeting the increasing requirements of their applications. D. Zhao (B) · Y. Yang · X. Song · Y. Liu · C. Sun · J. Zhou Shandong Provincial Key Laboratory for High Strength Lightweight Metallic Materials, Advanced Materials Institute, Qilu University of Technology (Shandong Academy of Sciences), Jinan 250014, China e-mail: [email protected] D. Zhao · Y. Yang Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China X. Song · J. Zhou Shandong Engineering Research Center for Lightweight Automobiles Magnesium Alloys, Advanced Materials Institute, Qilu University of Technology (Shandong Academy of Sciences), Jinan 250014, China © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_1

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During recent years, many studies have shown that the Mg–Zn–Al–Sn alloy has the potential to be a new kind of wrought magnesium alloy [2–5]. In our previous study, a ZAT422 alloy was successfully extruded at a temperature as low as 498 K and exhibited good mechanical properties. However, with the temperature increasing from 498 to 548 K, the yield strength of ZAT422 alloy decreased from 212 to 142 MPa and the elongation decreased from 27.6 to 23.1%, which was due to the grain size increasing with the extrusion temperature. Mn was proved to be effective in refining the dynamic recrystallization grain in magnesium alloy [6–9]. Therefore, a small amount of Mn (0.6 wt%) was added to ZAT422 alloy in the present study, which is aimed to improve the mechanical properties of the ZAT422 alloy extruded at a high temperature, and the microstructures and mechanical properties of the ZAT422 alloy and ZATM4220 alloy were studied.

2 Experimental Procedure The ZAT422 alloy and ZATM4220 alloy were prepared by semicontinuous casting in a protective atmosphere (0.5 vol% SF6 + 99.5 vol% CO2 ). The ingots were homogenized at 608 K for 4 h and 693 K for 4 h (T4) followed by water quenching at room temperature, and then machined to cylindrical-shaped samples with the size of Ø120 mm × 150 mm. Extrusion experiments were performed on an XJ-800SM extrusion machine. At first, the extrusion chamber and ingots were heated to 648 K for 2 h. Then, the ingots were extruded into bars with a diameter of 16 mm at an exit speed of 1–2 m min−1 and a reduction of 61. The extruded samples were aged at 70 °C for 24 h and 180 °C for 12 h followed by water quenching (T5). The samples for observation were cut out from the longitudinally central part of the extruded bars. The microstructure was examined by optical microscopy (OM) (Zeiss Axio Observer Alm), scanning electron microscopy (SEM) (JSM-840 instrument) equipped with electron backscatter diffraction (EBSD) (Oxford Instruments-hkl) detector, and a transmission electron microscopy (TEM) (Tecnai 20). The samples for EBSD analysis were additionally performed by electro-polishing in an ethanol–perchloric acid solution after mechanically polishing. HKL Channel 5 provided the orientation imaging microscopy (OIM) software. Tensile specimens with 25 mm in gage length and Ø5 mm in gage diameter, and cylindrical compression samples with the size of Ø8 mm × 12 mm were prepared for mechanical tests. The mechanical properties were tested on a universal testing machine (Instron AG-100kNG) at a crosshead speed of 1 mm min−1 .

Microstructures and Mechanical Properties of Extruded and Aged …

3

3 Results and Discussion 3.1 The Microstructures of the ZAT422 and ZATM4220 Alloys As shown in Fig. 1, both of the as-cast ZAT422 and ZATM4220 alloys exhibit a typical dendritic structure with α-Mg dendrite and semicontinuous second phase particles. And it can be seen that the grain sizes are similar to each other. Our previous study [10] has confirmed that the as-cast ZAT422 alloy consists of α-Mg, Mg32 (Al,Zn)49 and Mg2 Sn phases. After Mn was added into the ZAT422 alloy, the microstructure is shown in Fig. 1b, c. Besides the similar phases of ZAT422 alloy, some Al–Mn particles pointed in Fig. 1c can be observed in ZATM4220 alloys. Based on the energy-dispersive spectroscopy (EDS) spectrum (Fig. 1d) and literature [11, 12], those particles were identified as Al8 Mn5 . Figure 2 shows the microstructures and EDS spectrums of second phases in the asextruded alloys. As shown in Fig. 2a, b, the extrusion direction (ED) is pointed by the arrow direction, and it is found that completely dynamic recrystallization occurred in both of the two alloys. The average grain size of as-extruded ZAT422 is 19.1 μm. While the grain size of ZATM4220 decreases to 4.6 μm, and the microstructure

(a)

(b)

50μm

50μm

(c)

(d)

5μm

Fig. 1 Microstructure images (a, b, c) of the as-cast a ZAT422 alloy and b, c ZATM4220 alloy, d the EDS result of the particle pointed by the black arrow in c

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shows a zonal distribution along the extrusion direction. Therefore, it is concluded that Mn is effective in refining the dynamically recrystallized grains of ZAT422 alloy. As shown in Fig. 2c, d, some particles are observed in the microstructure of the two alloys. Based on the EDS results (Fig. 2e–g), most of the micro-scaled particles in ZAT422 alloy were confirmed to be Mg2 Sn, and a few nanoscaled particles with the size of 40–50 nm was observed in the interior of grain (Fig. 2h), which may be Mg2 Sn or MgZn precipitated during extrusion. The particles pointed in Fig. 2c, d were identified as Mg2 Sn, Al8 Mn5, and MgZn (Fig. 2e–g), and numbers of particles with the size less than 10 nm were observed in a grain of the as-extruded ZATM4220 alloy, which is considered to be α-Mn [13]. Figure 3 shows the EBSD results of the as-extruded ZAT422 alloy and ZATM4220 alloy. As depicted in Fig. 3a, b, orientation map is used as the background and low-angle grain boundaries (LAGBs) (2°–10°) are delineated by thin black lines, and high-angle grain boundaries (HAGBs) (10°–90°) are marked with coarse black lines. More LAGBs are observed within the DRXed grains interior of ZAT422 alloy than ZATM4220 alloy. And the misorientation distribution graphs (Fig. 3c, d) show the same trend, and it is found that there is a high-frequency peak near 86.3°, which ¯ tension twining in ZAT422 alloy. Some twins can be observed corresponds to {1012} in the grain pointed by the black arrows in Fig. 3a. This phenomenon agrees with our previous study [14], that is twin dynamic recrystallization (TDRX) playing an important role in the DRXed process of ZAT422 alloy. While, with the addition of 0.6 wt% Mn, the frequency of the LAGBs is decreased to 0.045, and the frequency peak near 86.3° is also lower than that of ZAT422. All the above show that the Mn addition in ZAT422 alloy enhances the recrystallization processing. Figure 3e, f shows the (0002) polar and ED inverse polar graphs of the two alloys. ¯ [0002] prismatic texture are observed in ZAT422. A (0002) basal texture and a [1010] The texture of ZATM4220 is similar with ZAT422, but the crystal direction of the ¯ to [21¯ 10]. ¯ And the max polar density basal texture is inclined about 20° from [1010] of ZAT422 alloy is 16.17, while, the ZATM4220 alloy is 5.77, meaning that the Mn addition weakened the texture of ZAT422.

3.2 The Mechanical Properties of the As-extruded and As-aged ZAT422 and ZATM4220 Alloys The mechanical properties of the as-extruded and as-aged ZAT422 and ZATM4220 alloys are listed in Table 1. Figure 4 shows the engineering stress–strain curves of ZAT422 and ZATM4220 alloys. It can be seen that ZATM4220 alloy exhibits better mechanical properties than ZAT422 alloy. With the addition of Mn, the tensile yield strength (TYS) of as-extruded ZATM4220 increases from 142.3 to 226.2 MPa, and the compression yield strength (CYS) increases from 139.0 to 229.9 MPa. The ultimate tensile strength (UTS) and ultimate compression strength (UCS) also increases by 5.9 and 7.7%, separately. As known from the microstructure observation, the grain

Microstructures and Mechanical Properties of Extruded and Aged …

(a)

5

(b)

ED

ED

20μm

(c)

20μm

(d) MgZn Mg2Sn Mg2Sn

Al 8Mn5

20μm

(e)

20μm

(f)

(g)

(i)

(h)

α-Mn

100nm

20nm

Fig. 2 Microstructure images (a, b, c, d, h, i) of as-extruded ZAT422 alloy (a, c, e, h) and ZATM4220 alloy (b, d, f, g, i), EDS results (e, f, g) of the particles pointed by the black arrow in c, d

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(a)

(b)

(c)

(d) 0.18

0.2

86°

0.18

Relative frequency

Relative frequency

0.2

0.1

0.0

0

10

20

30

40

50

60

70

80

86°

0.045

0.0

90

0

10

20

30

40

50

60

70

80

90

Misorientation angle (° )

Misorientation angle (° )

(e)

0.075

0.1

(f)

Fig. 3 EBSD results of the as-extruded ZAT422 alloy (a, c, e) and ZATM4220 alloy (b, d, f)

size of as-extruded ZATM4220 alloy is smaller than as-extruded ZAT422 alloy. Consequently, according to Hall Petch formula, the mechanical properties of the ZATM4220 alloy are improved. Besides, there are numbers of α-Mn particles with the size less than 10 nm in the as-extruded ZATM4220 alloy, which can hinder the movement of dislocations during the deformation process. However, the elongation (E) of the as-extruded ZATM4220 alloys increases little (from 23.1 to 23.8%). According to the microstructure of near tensile fracture surface of as-extruded ZAT422 and ZATM4220 alloys, as shown in Fig. 5, long cracks generated in the Al8 Mn5 particles are observed, which are more easily spread further, finally resulting in the fracture failure. After aging treatment, the tensile strength and compressive strength of both ZAT422 and ZATM4220 alloys are improved.

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Table 1 Tensile and compressive properties of extruded ZAT422 and ZATM4220 alloys Alloy

ZAT422

ZATM4220

Status

Tensile property

Compressive property

R

TYS (MPa)

UTS (MPa)

E (%)

CYS (MPa)

UCS (MPa)

Asextruded

142.3

307.2

23.1

139.0

446.9

0.97

As-aged

214.2

323.4

20.0

208.4

455.3

0.97

Asextruded

226.2

325.2

23.8

229.9

481.1

1.02

As-aged

270.2

336.9

20.0

261.8

475.2

0.96

Note R–tensile compression yield point asymmetry ratio

(b) As-extruded

500 400

ZTAM4220 ZTA422 ZTAM4220

300

ZTA422

200 Tensile curve Compressure curve

100 0 0.00

0.05

0.10

0.15

0.20

Engineering Strain (%)

0.25

Engineering Stress (MPa)

Engineering Stress (MPa)

(a)

As-aged

500 ZTAM4220

400

ZTA422 ZTAM4220

300

ZTA422

200 Tensile curve Compressure curve

100 0 0.00

0.05

0.10

0.15

0.20

0.25

Engineering Strain (%)

Fig. 4 Engineering stress–strain curves of ZAT422 and ZATM4220 alloys: a as-extruded, b as-aged

It is also seen from Table 1 and Fig. 4 that the two alloys exhibit good yield tension compression symmetry, and the tensile compression yield point asymmetry ratios (R) of the as-extruded ZAT422 and ZATM4220 alloys is 0.97 and 1.02, separately. According to the literature [5, 15], the existence of the texture of basal plane perpendicular to the extrusion direction can hinder the strain on c-axis direction ¯ 1011 ¯ twinning, which is favorable to obtain a higher compression caused by {1012} strength.

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(b)

(a)

Al 8 Mn5

10μm

10μm

Fig. 5 SEM images of near tensile fracture surface of as-extruded ZAT422 (a) and ZATM4220 (b) alloys

4 Conclusions (1) The addition of 0.6 wt% Mn has a negligible effect on the grain size of as-cast ZAT422 alloy, but α-Mn helps to decrease the dynamic recrystallization grain size of as-extruded ZAT422 alloy from 19.1 to 4.6 μm. ¯ [0002] prismatic texture are observed in both (2) A (0002) basal texture and a [1010] of the as-extruded ZAT422 alloy and ZATM4220 alloy. The Mn addition helps to reduce the max polar density and enhance the recrystallization processing of ZAT422 alloy. (3) Compared with ZAT422 alloy, the as-extruded and as-aged ZATM4220 alloy exhibits better mechanical properties. The TYS and CYS of the extruded ZATM4220 are 226.2 and 229.9 MPa, separately, exhibiting good yield tension compression symmetry. Both ZAT422 alloy and ZATM4220 alloy can be strengthened by aging treatment. Acknowledgements The work reported in this paper was supported by the National Key Research and Development Program of China (Project No. 2017YFB0103904), Shandong Province Key Research and Development Plan (Project No. 2017CXGC0404) and the Natural Science Foundation of Shandong Province (Project No. ZR2016EMB11).

References 1. B.L. Mordike, T. Ebert, Magnesium properties-applications-potential. Mater. Sci. Eng. A 302(1), 37–45 (2001) 2. S. Harosh, L. Miller, G. Levi, M. Bamberger, Microstructure and properties of Mg-5.6%Sn4.4%Zn-2.1%Al alloy. J. Mater. Sci. 42(24), 9983–9989 (2007) 3. T.T. Sasaki, K. Yamamoto, T. Honma, S. Kamado, K. Hono, A high-strength Mg-Sn-Zn-Al alloy extruded at low temperature. Scr. Mater. 59, 1111–1114 (2008) 4. W.L. Cheng, H.S. Kim, B.S. You, B.H. Koo, S.S. Park, Strength and ductility of novel Mg8Sn-1Al-1Zn alloys extruded at different speeds. Mater. Lett. 65(11), 1525–1527 (2011)

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5. D.Q. Zhao, J.X. Zhou, Y.T. Liu, X.G. Dong, J. Wang, Y.S. Yang, Microstructure and mechanical properties of Mg-4Zn-2Al-2Sn alloys extruded at low temperatures. Acta Metall. Sin. 50(1), 41–48 (2014) 6. M. Celikin, A.A. Kaya, M. Pekguleryuz, Effect of manganese on the creep behavior of magnesium and the role of α-Mn precipitation during creep. Mater. Sci. Eng. A 534(534), 129–141 (2012) 7. P. Cao, M. Qian, D.H. Stjohn, Effect of manganese on grain refinement of Mg-Al based alloys. Scr. Mater. 54(11), 1853–1858 (2006) 8. F. Zarandi, G. Seale, R. Verma, E. Essadiqi, S. Yue, Effect of Al and Mn additions on rolling and deformation behavior of AZ series magnesium alloys. Mater. Sci. Eng. A 496(1–2), 159–168 (2008) 9. S. Wang, R. Ma, L. Yang, Y. Wang, Y. Wang, Precipitates effect on microstructure of asdeformed and as-annealed AZ41 magnesium alloys by adding Mn and Ca. J. Mater. Sci. 46(9), 3060–3065 (2011) 10. D.Q. Zhao, X.G. Dong, X.E. Zhang, A.J. Gao, J.X. Zhou, Y.S. Yang, Microstructure and tensile properties of Mg-4Zn-2Sn-2Al alloy. Mater. Sci. Forum 747–748, 398–403 (2013) 11. S.S. Lun, D. Dube, R. Tremblay, Characterization of α-Mn particles in AZ91D investment castings. Mater. Charact. 58(10), 989–996 (2007) 12. Y.M. Kim, C.D. Yim, B.S. You, Grain refining mechanism in Mg-Al base alloys with carbon addition. Scr. Mater. 57(8), 691–694 (2007) 13. X.Y. Fang, D.Q. Yi, J.F. Nie, X.J. Zhang, B. Wang, L.R. Xiao, Effect of Zr, Mn and Sc additions on the grain size of Mg-Gd alloy. Alloys Compd. 470(1–2), 311–316 (2009) 14. D.Q. Zhao, Y.S. Yang, J.X. Zhou, Y. Liu, C.W. Zhan, Dynamic microstructural evolution in mg-4Zn-2Al-2Sn alloy during hot deformation. Mater. Sci. Eng. A 657, 393–398 (2016) 15. Y.N. Wang, J.C. Huang, The role of twinning and untwinning in yielding behavior in hotextruded Mg-Al-Zn alloy. Acta Mater. 55(3), 897–905 (2007)

Microstructural Evolution of an Al–Zn–Mg–Cu Aluminum Alloy During an Optimized Two-Step Homogenization Treatment Hongwei Yan, Xiwu Li, Zhihui Li, Shuhui Huang, Hongwei Liu, Lizhen Yan, Wen Kai, Yong’an Zhang and Baiqing Xiong Abstract The microstructural evolution of an Al–Zn–Mg–Cu aluminum alloy during an optimized two-step homogenization treatment was investigated by optical microscopy (OM) and scanning electron microscopy (SEM) equipped with energydispersive spectrometer (EDS) and differential scanning calorimetry (DSC). It mainly focused on secondary phase transformation and dissolution. Methods such as and differential scanning calorimetry (DSC) was used to investigate the heat change of each specimen as an indication of phase transformation involved in the homogenization. The results showed that the lamellar eutectic phase had the trend of spheroidizing during the treatment at 400 °C, and the T-Al2 Mg2 Zn3 phase was transformed into S-Al2 CuMg phase. Further dissolution of T phase was not obvious at second step homogenization even after the holding time was extended to more than 24 h.

1 Introduction Al–Zn–Mg–Cu alloys are widely applied in aerospace, aviation, and transportation industry, because of its outstanding mechanical properties and corrosion resistance. 7xxx aluminum alloy has the highest strength within all the commercial aluminum alloys, in that case, these pre-stretched plates have been used in the manufacture of airliners such as Airbus A380 and Boeing 777. However, microstructure defects such as dendrite, segregation, coarse eutectic structure at grain boundaries and uneven distribution of alloy elements are more prone to appear during the casting process because of the design concept of high alloying degree. Without proper treatment in the processes followed, these defects might result in the deterioration of key properties such as strength, plasticity, corrosion resistance, etc. [1–3]. H. Yan (B) · X. Li · Z. Li · S. Huang · H. Liu · L. Yan · W. Kai · Y. Zhang · B. Xiong State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co., Ltd., Beijing 101407, China e-mail: [email protected] B. Xiong GRINM Group Co., Ltd., Beijing 100088, China © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_2

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Table 1 Chemical composition of the studied 7055 aluminum alloy (wt%) Element

Cu

Mg

Zn

Zr

Fe

Si

Al

Content

2.21

2.10

7.98

0.096

0.042

0.014

Bal.

In the development of aluminum alloy industry, both theory and practice have proved that homogenization plays an essential role in eliminating segregation and resolving the low melting point eutectic phases. For aluminum alloys containing zirconium, homogenization can also promote the dispersion of Al3 Zr particles. Finely dispersed Al3 Zr particles inhibit recrystallization by pinning grain boundaries from migrating, thus preventing the grain growth that reduces strength [4, 5]. Therefore, homogenization of ingots is a fundamental process of the manufacture of ultrahigh strength aluminum alloys, and it has great significance for obtaining highperformance aluminum materials. It is generally acknowledged that the second phased in as-cast structure of 7055 aluminum alloy are mainly composed of MgZn2 (η phase), T phase, S phase and Fe-rich Al7 Cu2 Fe phase [6–9]. During homogenization, the phases formed during the nonequilibrium solidification of the casting process gradually dissolve, and segregation is gradually eliminated, but the change of Fe-rich phase is not obvious [10]. The investigation of detailed microstructure transformation of the eutectic structure based on published literature still lacks in-depth explanation. In this paper, the microstructural evolution of 7055 aluminum alloy ingot during an optimized twostep homogenization treatment was studied by DSC and SEM, and it is focused on the evolution of the eutectic phase [11].

2 Experimental Procedure The direct chill (DC) casting ingots were provided by Northeast Light Alloy Co., Ltd. and the chemical composition of the ingots was shown in Table 1. Small specimens with the size of 10 × 10 × 10 mm were cut out of the ingot, and the heat treatment was carried in an air circulation box furnace. Based on the former study and factory production, a two-step homogenization was chosen, which is 400 °C/10 h + 470 °C/t h (t  0, 24, 32, 40, 48, 60, 72). The interval between the first and the second stage was 3 h. The specimens were labeled from 7055-1 to 7055-7 corresponding to the increase of holding time of the second stage. After homogenization, the specimens were quenched into water to retain their structures. The microstructures were observed by Axiovert 200 MAT optical microscope (OM) and JEOL JSM 7001F field emission scanning electron microscope (SEM) and the phase dissolution was investigated by differential scanning calorimeter (DSC). The OM specimens were polished and etched by Keller reagent, and the SEM specimens were electropolished. DSC tests were performed in NETZSCH DSC 404F3A, and the scanning rate is 10 K/min.

Microstructural Evolution of an Al–Zn–Mg–Cu Aluminum Alloy …

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3 Results and Discussion Figure 1 shows the microstructure of DC casting ingot. The investigated ingot shows typical casting structure, and it mainly consists of α-Al solid solution and the layered low-melting eutectic phase formed at the grain boundary. And a small amount of precipitates formed during solidification were observed in the grains. Some grey areas were observed in the eutectic structure, as indicated by the arrow in Fig. 2. According to the EDS results (Table 2), the second phases marked by A and B are Fe-rich phase and Al2 CuMg, respectively. Meanwhile, intragranular and coarse MgZn2 precipitates were observed near the grain boundary [10, 12]. The evolution of second phases dissolution was studied by observing the microstructure of specimens after different holding time. As shown in Fig. 3, when the first step heat treatment at 400 °C finished and the specimen was just heated up to 470 °C, the original eutectic structure at the grain boundary is still observable (Fig. 3a). The layered quaternary phase tends to spheroidize and dissolve, and the interface between different layers begins to become blurred as shown in Fig. 4. When the holding time of the second stage was extended to 24 h (Fig. 3b), the layered

Fig. 1 The microstructure of studied aluminum alloy DC casting ingot a OM; b SEM backscattered electron image Fig. 2 SEM backscattered electron image of the layered low-melting eutectic phase

14 Table 2 EDS analysis of phase composition (atomic ratio)

H. Yan et al.

Position

Element Cu

Fe

A

74.71

1.95

1.66

11.04

10.65

B

74.94

4.18

10.75

10.13

–

C

79.55

3.38

9.59

7.48

–

D

33.58

18.93

30.37

17.11

–

E

58.73

2.03

20.58

20.85

–

Al

Zn

Mg

Fig. 3 SEM backscattered electron image of specimens with various homogenization processes a 400 °C/10 h + 470 °C/0 h; b 400 °C/10 h + 470 °C/24 h; c 400 °C/10 h + 470 °C/32 h; d 400 °C/10 h + 470 °C/40 h; e 400 °C/10 h + 470 °C/48 h; f 400 °C/10 h + 470 °C/60 h

quaternary phase has completely dissolved, and the residual AlZnMgCu quaternary phase shows an obvious feature of spheroidization. During the diffusion of solute elements, the interface between the second phase and the matrix gradually becomes smoother, and the original continuous quaternary phase gradually becomes disconnected. The residual phase is mainly composed of AlZnMgCu quaternary phase, a small number of S phase with large size and iron-rich phase. As the holding time was further extended, the dissolution of the residual phases was not obvious. Neither the shape nor the species showed distinct changes, as shown in Fig. 3c–f. Furthermore, the phase transformation from T phase to S phase was observed during the first stage of homogenization. As shown in Fig. 5, second phases with grey contrast are found near the quaternary phase in the specimen after the first stage of homogenization. These phases are only a few microns in size, and closely adjacent to the quaternary phase. EDS analysis of them shows that these particles mainly contain Al, Cu, and Mg, as listed in Table 2. Therefore, this phase is presumed to be S phase, but the morphology is different from the S phase observed in the ascast specimen which is mainly distributed along the quaternary phase at the grain

Microstructural Evolution of an Al–Zn–Mg–Cu Aluminum Alloy …

15

Fig. 4 Backscattered electron SEM image of specimens before and after first stage homogenization a as-cast; b after heat treatment of 400 °C/10 h

Fig. 5 The residual phases in the sample after first stage of homogenization at 400 °C/10 h

boundary, as marked by B in Fig. 2. While the quaternary phase was observed to be surrounded by the S phase. Hence, the transformation from AlZnMgCu quaternary phase to S phase happened during homogenization because Zn atom defused into the matrix preferentially. And furthermore, this transformation occurred during the first stage of homogenization at 400 °C. The microstructure transformation during the homogenization process mainly involves the re-dissolution and precipitation of the second phase, so it is inevitably accompanied by the endothermic or exothermic reaction during the heating process. Considered that microstructure characterization is relatively local, DSC experiments were carried out in this study. The position and peak height information of the heat reaction peaks were obtained to study the evolution of the second phases during the homogenization process. As is shown in Fig. 6a, a series of DSC curves of samples with different homogenization times were obtained. A strong endothermic peak appears at 469.5 °C, which represents the melting of the non-equilibrium solidified eutectic phase in the as-cast microstructure. After the first step of homogenization, this endothermic peak is still obvious, but its intensity is weakened. When the second stage homogenization is

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Fig. 6 DSC analysis of specimens with various homogenization processes a DSC curves; b the onset temperature of eutectic phase melting peak

carried out for more than 24 h, a significant reduction of the endothermic peak intensity is observed, and a new endothermic peak appears at 485 °C. As is shown in Fig. 6b, the onset points of eutectic phase melting peaks of all eight samples are compared. After the first step of homogenization, the onset temperature increases by 5 °C. As the holding time of the second step is further extended, the onset temperature is increasing gradually, and it is stable at 475 °C when the holding time exceeds 40 h. For all specimens with the second step holding time more than 24 h, another endothermic peak is observed at 485 °C. The onset temperature of the endothermic peak which indicates the melting of the quaternary phase was increased from 469.5 to over 474 °C, and that ensure the ingot would not be over burnt during the second step of homogenization at 470 °C. During the homogenization process, the primary driving force for the residual phase to redissolve is the concentration difference across the phase boundary. During the first step of homogenization, even though the redissolution of residual phase at a large scale cannot happen because of insufficient activation energy, the diffusion of solute atoms stabilizes the quaternary phase and increases its melting temperature. When the holding time at 470 °C was over 24 h, the significant drop of the endothermic peak intensity indicates that most of the quaternary phase has redissolved. The endothermic peak at 485 °C was not observed on the DSC curves of as-cast specimen and first step homogenized specimen. That is because the volume fraction of S phase in these specimens are relatively low and the amount of heat change is not enough to form an endothermic peak during the scanning process in a DSC. And due to the transformation from the quaternary phase to S phase occurred during the second step of homogenization, the volume fraction of S phase increase and the endothermic peak appears. And this is consistent with the microstructure observation.

Microstructural Evolution of an Al–Zn–Mg–Cu Aluminum Alloy …

17

4 Conclusions (1) The two-step homogenization of 400 °C/10 h + 470 °C/72 h showed a great effect for the heat treatment of as-cast Al–Zn–Mg–Cu alloy, but a complete dissolution of eutectic phase was still not achieved. When the treatment time of the second step homogenization exceeded 24 h, no significant improvement of the microstructure was observed. (2) During the first step of homogenization process (400 °C/10 h), the evidence of transformation from Al–Zn–Mg–Cu quaternary phase to S phase was observed, and the lamellar eutectic phase showed the trend of spheroidizing. Acknowledgements This work was supported by the National Key R&D Program of China (No. 2016YFB0300803).

References 1. F. Xie, X. Yan, L. Ding, F. Zhang, S. Chen, G. Men, Y. Chang, Mater. Sci. Eng. A 355(1–2), 144–153 (2003) 2. X.M. Li, M.J. Starink, Mater. Sci. Technol. 17(11), 1324–1328 (2001) 3. Y. Deng, Z.M. Yin, F.G. Cong, Intermetallics 26, 114–121 (2012) 4. J.D. Robson, P.B. Prangnell, Acta Mater. 49(4), 599–613 (2001) 5. Z. Guo, G. Zhao, X.G. Chen, Mater. Charact. 102, 122–130 (2015) 6. C. Mondal, A.K. Mukhopadhyay, Mater. Sci. Eng. A 391(1–2), 367–376 (2005) 7. U. Tenzler, E. Cyrener, G. Tempus, Aluminium 75(6), 524–530 (1999) 8. L.L. Rokhlin, T.V. Dobatkina, N.R. Bochvar, E.V. Lysova, J. Alloy. Compd. 367(1–2), 10–16 (2004) 9. S.D. Liu, Y.B. Yuan, C.B. Li, J.H. You, X.M. Zhang, Met. Mater. Int. 18(4), 679–683 (2012) 10. N.K. Li, J.Z. Cui, Trans. Nonferrous Met. Soc. China 18(4), 769–773 (2008) 11. K. Chen, H. Liu, Z. Zhang, S. Li, I.T. Richard, J. Mater. Process. Technol. 142(1), 190–196 (2003) 12. F.G. Cong, G. Zhao, F. Jiang, N. Tian, R.F. Li, Trans. Nonferrous Met. Soc. China 25(4), 1027–1034 (2015)

Microstructural Evolution and Phase Transformation of Al–Mg–Si Alloy Containing 3% Li During Homogenization Xiaokun Yang, Baiqing Xiong, Xiwu Li, Lizhen Yan, Zhihui Li, Yong’an Zhang, Hongwei Liu, Shuhui Huang, Hongwei Yan and Kai Wen Abstract The microstructural evolution and phase transformation of cast Al–1.5Mg–0.6Si–3Li (mass %) alloy during homogenization were investigated. The results show that severe dendritic segregation exists in the as-cast ingot. Mg and Si elements segregate at grain boundaries to form intermetallic Mg2 Si phase. In addition, there also exists Li-containing phases, including T-Al2 LiMg and δ-AlLi phase in the α-Al matrix. These Li-containing phases completely dissolve into the matrix and the segregation of dendrite is eliminated after homogenization at 570 °C for 24 h. During the homogenization, most of the Mg2 Si phase at grain boundaries disappear, but AlLiSi ternary compounds precipitate and disperse at interior and boundary of grains because of the strong binding capacity between Li and Si element. The AlLiSi phase is detrimental to the properties of the alloy, therefore homogenization treatment may be not profitable for microstructural refinement of Al–1.5Mg–0.6Si–3Li alloy.

1 Introduction The demand for lightweight structure material is increasing for the aerospace industry. A reduction of total mass for airplanes can improve their payload capability [1, 2]. Al–Li alloy has superior properties, such as low density, high specific strength and stiffness, and high damage resistance because of the contribution of Li addition [3–5]. According to the previous studies, the addition of 1 wt% Li into Al alloy can approximately decrease 3% density and increase 6% elastic modulus [6]. Therefore, it has gained attention from their employees in aircraft applications since Al–Li alloy was invented .

X. Yang · B. Xiong (B) · X. Li · L. Yan · Z. Li · Y. Zhang · H. Liu · S. Huang · H. Yan · K. Wen State Key Laboratory of Nonferrous Metals and Processes, GRINM Group Co. Ltd., Beijing 101407, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_3

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The third-generation Al–Li alloy has been widely applied at present. It is mainly Al–Cu–Li–Mg series alloy with minor Zr, Ag, Zn, and other elements, and the base alloy is 2xxx series. However, the proportion of Cu increase as the main alloying element results in the effect of weight loss declining for Al–Li alloy. The density of this generation Al–Li alloy is about 2.7 g/cm3 , which is higher than the past two generations in general [7]. Al–Mg–Si series alloy or 6xxx series alloy is extremely useful commercial medium-strength aluminum alloy and has been widely used in automobile and aerospace industry [8–10]. Addition of Li to Al–Mg–Si alloy promotes the formation of δ -Al3 Li and β -Mg2 Si dual phases strengthening so as to achieve a novel low-density Al–Li alloy, which is the base alloy is 6xxx series. Some researchers have carried out experiments to study the influence of Li addition on Al–Mg–Si alloy [11–14], but their work mainly focused on the characterization of precipitation during aging treatment and there was little research on the microstructural evolution of Al–Mg–Si–Li alloy during homogenization. The result of homogenization has an important effect on the microstructure, heat treatment process and mechanical properties of the alloy. Therefore, the aim of the present research is to investigate the influence of Li addition on microstructure features of Al–Mg–Si alloy during homogenization.

2 Experimental The composition of the experimental material is shown in Table 1. The alloy was melted in a medium frequency induction furnace in a graphite crucible. First, pure aluminum was loaded in a preheated crucible. After this pure aluminum was melted, Al–12 wt% Si master alloys and pure Mg were added to the melt at 700 °C. Second, pure Li was covered with aluminum foil and plunged into the melt by a metal bell jar. After a few minutes, AlTiB (94:5:1) was added in the melt as a grain refiner. High purity Ar gas and LiCl–LiF flux cover were used to protect melt against air. The molten alloy was poured into a cast iron mold under Ar gas protection. The mold cavity size was approximately 300 mm × 200 mm × 50 mm. After cooling, the experimental samples were cut 10 × 10 × 15 mm by linear cutting machine from the center of the as-cast ingot for further heat treatment. The homogenization treatment was performed at 570 °C for 24 h. Differential scanning calorimetry (DSC) tests of the as-cast and homogenized samples were carried out by NETZSCH DSC 404 F3 instrument in an argon atmosphere from 25 to 650 °C with a heating rate of 10 °C/min. The size of the samples was 4 × 1.2 mm. The reference standard samples were pure aluminum with a similar shape and mass. The metallographic samples were milled by SiC water polishing

Table 1 Chemical composition of the experimental alloy (mass %)

Li

Mg

Si

Fe

Al

3.00

1.50

0.64

15° (%)

15

H112

69.3

11.2

19.5

60

H112

81.9

12.3

5.8

Study on Stabilization Treatment of Al–Mg Alloy 5E83-H112

33

angle is 2°–5°, the green lines represent a grain boundary whose misorientation angle is 5°–15°, and the black lines represent a high-angle grain boundary(>15°). Table 3 shows the statistical results of percentages of different misorientation in the same area. From the table, it can be seen that the grain boundaries of the two alloys are mainly dominated by low-angle grain boundaries. In the same area, the number of grain boundary in Fig. 2a is significantly more than that in Fig. 2b, so the number of low-angle grain boundary in Fig. 2a is much larger than that in Fig. 2b. The formation of subgrain boundaries inside the grains is due to the dynamic recovery and partial recrystallization during the hot rolling process, the dislocations were eliminated by thermal activation, the cell walls became sharp, the dislocation density decreased, and the entanglement became regular, subgrain and subgrain boundaries were formed.

3.3 Intergranular Corrosion and Exfoliation Corrosion Table 4 shows the weight loss of the alloys in different states, test specimens were sensitized by aging at 100 °C for 168 h. The weight loss value per unit area of the two hot-rolled sheets before sensitization are 3.1 and 2.2 mg/cm2 , suggesting that the alloy has good resistance to intergranular corrosion. After sensitization treatment, the weight loss value of the 15 mm hot-rolled sheet increased to 40.8 mg/cm2 , which has entered the severe corrosion zone; the weight loss value of the 60 mm hot-rolled sheet increased to 17.6 mg/cm2 , which is in the passivation zone. In order to further study the intergranular corrosion of the passivation area, it needs to be observed by metallography. As shown in Fig. 3, the alloy does not undergo significant lamellar corrosion, indicating that the alloy has good resistance to intergranular corrosion. Figure 4 is a macroscopic photograph of the exfoliation corrosion in hot-rolled state and sensitized state of the two alloys. The surface of the hot-rolled alloy has no appreciable attack and no peeling. The surface of the sensitized alloy has slight pitting corrosion, which indicates that the alloy has good resistance to exfoliation corrosion. The corrosion levels corresponding to the respective states are listed in Table 5. From the corrosion results, it can be seen that the intergranular corrosion resistance of hot-rolled sheet which is 15 mm is relatively poor. D’Antuono et al. [7] found that in many instances, the β phase formation preferred low-angle rather than high-angle grain boundaries. Low-angle boundaries are susceptible to sensitization. Therefore,

Table 4 Weight loss of intergranular corrosion Thickness (mm)

Condition

Weight loss per unit area before sensitization (mg/cm2 )

Weight loss per unit area after sensitization (mg/cm2 )

15

H112

3.1

40.8

60

H112

2.2

17.6

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the intergranular corrosion performance of 15 mm hot-rolled sheet is poor may be due to lots of low-angle grain boundaries in the alloy. However, the rate of exfoliation corrosion of two alloy plates is the same, it indicates that the alloy is insensitive to exfoliation corrosion but rather to intergranular corrosion. Therefore, efforts should be made to improve intergranular corrosion resistance.

3.4 Stabilization Treatment It can be seen from the above results that the 15 mm hot-rolled sheet has been severely corroded after accelerated corrosion, meaning that it has poor resistance to intergranular corrosion. Therefore, in order to ensure good IGC resistance of the alloy during long-term use, it is necessary to perform a stabilization annealing treatment. Figure 5 shows the change of weight loss of the 15 mm alloy after annealing at 240,

Fig. 3 Intergranular corrosion depth of 60 mm hot-rolled sheet

(a) 15mm -H112

(b) 15mm-sensitized

(c) 60mm-H112

Fig. 4 Exfoliation corrosion morphology of two hot-rolled sheets

(d) 60mm-sensitized

Study on Stabilization Treatment of Al–Mg Alloy 5E83-H112 Table 5 The exfoliation corrosion levels

35

Thickness (mm)

Condition

Corrosion morphology

Rating

15

H112

Rolling surface is smooth, no appreciable attack

N

15

Sensitized

Rolling surface is smooth, a small number of pitting

PA

60

H112

Rolling surface is smooth, no appreciable attack

N

60

Sensitized

Rolling surface is smooth, a small number of pitting

PA

250, 260 °C for 1, 2, 4, 8, 12, 16, 20, 24 h. It can be seen that annealing at 240 °C, with the extension of annealing time, the weight loss curve quickly enters the sensitive area from the insensitive area. Annealing at 250 °C, the curve enters the medium sensitive area from the insensitive area and returns to the insensitive area after annealing for 16 h, however, after the sensitization treatment at 100 °C/168 h, the curve shows a downward trend but it is still in a severely corroded area. Annealing at 260 °C, the curve enters the insensitive area from the medium sensitive area after annealing for 4 h. From the weight loss curve after sensitization, it reaches the medium sensitization zone after annealing for more than 16 h and has been showing a downward trend. Figure 6 shows the intergranular corrosion depth after annealing at 260 °C for 20 and 24 h, it can be observed that the alloy has no obvious lamellar corrosion, indicating that the alloy has good resistance to intergranular corrosion. Figure 7a shows the hardness values of the alloy after annealing at different temperatures for different times. The hardness of the annealing state is lower than that of the initial state of H112, but it does not change significantly with annealing time. After sensitization at 100 °C/168 h, the hardness value did not decrease significantly compared to before sensitization, and the alloy showed a good stability of mechanical property. Figure 7b is a conductivity test of the alloy before and after sensitization at different temperatures and time. It can be seen that the conductivity increases slightly with the extension of annealing time at the same temperature, but it has no obvious change after sensitization. It is indicated that the β phase in the alloy after annealing will precipitate along the grain boundary, but the increase of conductivity is not obvious maybe due to the limited amount of supersaturated Mg element in the alloy.

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3.5 TEM Microstructure Figure 8 is the TEM photograph of the alloy of 15 mm. It can be seen from Fig. 8a that the alloy grain can be elongated after large thermal deformation, and the dislocation density is low, which is due to the high temperature during hot rolling. It would cause alloy to recover, and form many subgrains inside the deformed grains. The subgrain boundaries are clearly visible, but there are no large changes such as dissociation and migration, there are still dislocation entanglements in the structure. As shown in Fig. 8b, there are many long rod-shaped Al6 Mn particles and spherical Al3 (Er, Zr) particles dispersed in the alloy. These particles are distributed in the crystal and pinned on the grain boundaries or dislocations, they can be used to hinder the movement of dislocations and the subgrain boundaries, and form a polygonal stabilized

Fig. 5 Weight loss curve at different annealing temperatures for different time

Fig. 6 Intergranular corrosion depth after annealing at 260 °C for 20 and 24 h

Study on Stabilization Treatment of Al–Mg Alloy 5E83-H112

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Fig. 7 a Hardness curves of 5E83 alloy at different conditions; b conductivity curves of 5E83 alloy at different conditions

subgrain structure at high deformation temperature to improve the strength of the alloy [8]. Also, the precipitation of these second phases inhibits recrystallization while increasing the dynamic recrystallization temperature. Figure 8c is the grain boundary of the alloy microstructure. No precipitation of the β phase is observed, this is because the high temperature during hot rolling makes the β phase dissolve in the matrix, therefore, the alloy exhibits good resistance of intergranular corrosion performance. Figure 9 is the TEM image of the 15 mm hot-rolled sheet which is sensitized and stabilized. Figure 9a shows the deformed structure after sensitization, the grain still remains fibrous structure. From Fig. 9b, it can be observed that there is a continuous β phase on the grain boundary. This is because the β phase is easier to nucleate at a grain boundary or a region with higher energy such as the interface between the second phase and the matrix [9]. Figure 9c is the deformation recovery microstructure of the alloy treated at 260 °C/24 h + 100 °C/168 h. The subgrain is multi-lateralized inside, and the grain boundary migrates to form the subgrain; Fig. 9d is the grain boundary, the β phase can be observed to be intermittently distributed along the grain boundary.

Fig. 8 TEM micrographs of 5E83 alloy-H112 (15 mm): a deformed microstructure; b second phase particle; c grain boundary

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Fig. 9 TEM micrographs of 5E83 alloy (15 mm) sensitized and stabilized: a deformed microstructure at 100 °C/168 h; b β phase on grain boundary at 100 °C/168 h; c deformed microstructure at 260 °C/24 h + 100 °C/168 h; d β phase on grain boundary at 260 °C/24 h + 100 °C/168 h

The intergranular corrosion performance of Al–Mg alloy is greatly affected by the distribution of β phase. Since the corrosion potential of the β phase is higher than that of the aluminum matrix, the alloy is sensitive to IGC when the β phase is continuously distributed on the grain boundary [10]. Due to the grains of the hot-rolled sheet are elongated after being rolled, the β phase will be uniformly precipitated along the continuous grain boundary to form a β-phase network film after sensitization, thereby causing the intergranular corrosion resistance of the alloy to decrease. The reason of the β phase is intermittently precipitated along the grain boundary, is that the β phase is precipitated firstly from the matrix after annealing at 260 °C, and then coarsened until to fracture, so it is not continuously distributed on the grain boundary. This improves the resistance to intergranular corrosion of the alloy.

4 Conclusions Based on the above analysis, the following conclusions are drawn: (1) It can be seen from the EBSD structure of the two alloys that the interface of the two alloys is dominated by low-angle grain boundaries, while the intergranular

Study on Stabilization Treatment of Al–Mg Alloy 5E83-H112

39

corrosion performance of the 15 mm hot-rolled sheet is poor because of the low-angle grain boundaries of the 15 mm plate are more than the 60 mm. (2) The alloy was recovered and recrystallized after annealing at 240–260 °C, which can result in the strength decrease. (3) As the annealing temperature and time increases, the β phase precipitates along the grain boundary and the electrical conductivity increases. (4) After the sensitization of the hot-rolled sheet, the β phase is easily distributed on the grain boundary, and the intergranular corrosion performance of the sheet is bad; after treated at 260 °C/ 20 h or more than 20 h, there is only a small amount or no β phase on the grain boundary after sensitization, and the intergranular corrosion performance is good. Acknowledgements The authors are pleased to acknowledge the financial support received from the following projects (in no particular order). The National Key Research and Development Program of China (2016YFB0300804 and 2016YFB0300801), and the National Natural Science Fund for Innovative Research Groups (Grant No. 51621003). The Construction Project for National Engineering Laboratory for Industrial Big-data Application Technology (312000522303). National Natural Science Foundation of China (No. 51671005 and 51701006), Beijing Natural Science Foundation (2162006) and Program on Jiangsu Key Laboratory for Clad Materials (BM2014006).

References 1. C. Lavender, M. Smith, Aluminum Product Applications in Transportation and Industry, vol. 29 (ASME, Center for Research and Technology Development, 1994), p. 123 2. I.J. Polmear, Light Alloys: Metallurgy of the Light Metals (Edward Arnold, a division of Hodder Headline PLC, London, 1995) 3. S. Toros, F. Ozturk, I. Kacar, Review of warm forming of aluminum–magnesium alloys. J. Mater. Process. Technol. 207(1), 1–12 (2008) 4. L.P. Meng, Z.H. Bai, B.H. Luo, Effects of Mg, Sc and annealing temperature on Al-Mg-Sc alloy. Min. Metall. Eng. (2003) 5. Braun, Effect of thermal exposure on the corrosion properties of an Al-Mg-Sc alloys sheet. Mater. Sci. Forum 331:II(4), 9–14 (2000) 6. S.P. Wen, K.Y. Gao, Y. Li, H. Huang, Z.R. Nie, Synergetic effect of Er and Zr on the precipitation hardening of Al-Er-Zr alloy. Scr. Mater. 65, 592–595 (2011) 7. D.S. D’Antuono, J. Gaies, W. Golumbfskie et al., Grain boundary misorientation dependence of β phase precipitation in an Al-Mg alloy. Scr. Mater. 76, 81–84 (2014) 8. K.L. Kendig, D.B. Miracle, Strengthening mechanisms of an Al-Mg-Sc-Zr alloy. Acta Mater. 50, 4165–4169 (2002) 9. J. Yan, A.M. Hodge, Study of β precipitation and layer structure formation in Al 5083: the role of dispersoids and grain boundaries. J. Alloy. Compd. 703, 242–250 (2017) 10. H. Yukawa, Y. Murata, M. Morinaga et al., Heterogeneous distributions of magnesium atoms near the precipitate in Al-Mg based alloys. Acta Metall. Mater. 43(2), 681–688 (1995)

Simulation and Experimental Study on Hot Forging Process of SiCp/2A14 Composite Hansen Zheng, Zhifeng Zhang and Yuelong Bai

Abstract Taking the trackboard as the research object, the hot forging process of SiCp/2A14 composite slug prepared by stirring and casting method was studied by using numerical simulation and experimental research. The effects of slug size, hot forging pressure, hot forging time, and the heat treatment process on the formability and mechanical properties of the product were studied. The results show that the size of the blank of the similar volume has a great influence on the formability of the part. If the thickness is too thick or too thin, it will lead to die filling dissatisfaction; the higher the hot forging pressure, the more times of forging, the better the formability; after solid solution at 500 °C for 6 h, the parts aging for 8 and 12 h can obtain satisfactory mechanical properties, but it is found that as the aging time increases, the tensile strength of the composite increases while the elongation decreases.

1 Introduction The silicon carbide particle-reinforced aluminum matrix composites, characterized by the advantages of high specific strength, high specific rigidity, good wear resistance, and low coefficient of thermal expansion, have wide application prospects in the fields of defense and military, rail transportation, aerospace, and electrical appliances [1–3]. Although the composite parts directly formed by the stirring and melting method are low in cost, their plasticity is often low, and they need to be reinforced by thermal deformation to meet the user requirements. It is well known that it is difficult to obtain Al matrix composite parts with integrity by adopting the conventional deformation process, and new hot forging process need be developed. With the rapid development of computers and computing technologies, numerical simulation methods represented by the finite element method have been widely used in metal forming process analysis [4], which can make it possible to be low cost in research and development. H. Zheng · Z. Zhang (B) · Y. Bai General Research Institute for Nonferrous Metals, Beijing, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_5

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In this paper, trackboard was taken as the research object, and numerical simulation was used as a guide. Combined with experimental production, the hot forging process of SiCp/2A14 aluminum matrix composites prepared by stirring casting method [5] was performed, and the effects of slug size, hot forging pressure, hot forging time, and the heat treatment process on the formability and mechanical properties of the product were studied. The results are expected to provide certain guiding significance for the production process of similar products.

2 Experimental Section 2.1 Numerical Simulation Numerical simulation can greatly improve work efficiency, shorten cycle and reduce cost, and provide reference and guidance for subsequent experimental processing. The hot forging model of the trackboard was divided into three parts, and the solid modeling was performed in the UG drawing software, as shown in Fig. 1. The threedimensional schematic diagram of the die after assembly is shown in Fig. 2. The part model was imported in the Deform-3D software preprocessing module. Because the top and bottom die were not considered the force and deformation during the forging process, the top and bottom die were set as rigid, and the slug was set as plastic. After assembling, the initial speed of the bottom die and the workpiece was set to 0 according to the requirements of the experiment. The top die was the active part and the slug was the follower. The absolute method was adopted to divide the grid. Since only the stress distribution and forming conditions of the blank was considered in this experiment, only the blanks are meshed. The set mesh was a tetrahedral mesh with a size ratio of 2 and a minimum size of 1.5 mm. The calculation step was 1 mm. After the simulation was completed, the post-processing module in the Deform3D software was used to analyze the result. In this paper, three sets of different

(a) bottom die

(b) top die

Fig. 1 3D drawing model of hot-forged SiC/2A14 composite trackboard

(c) slug

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Fig. 2 Schematic diagram of three-dimensional die after assembly Table 1 Simulation process parameters

Number

Size of slug (mm)

Pressure (t)

1

120 × 120 × 41

400

2

120 × 120 × 41

800

3

110 × 110 × 50

800

forging conditions were used to simulate the data. The simulation process plan is listed in Table 1. In this experiment, the die temperature is set to 300 °C, the blank temperature is set to 480 °C, the pressing speed is set to 25 mm/s, and the contact relationship of slug is secondary. The friction type is shear friction, and the friction coefficient is 0.3, the interface heat transfer coefficient is 11, and the rest is 5.

2.2 Experimental Process The process parameters obtained by numerical simulation still need to be verified and corrected through specific experiments. The slug was heated in a box resistance furnace and the heating speed was relatively slow, but heating is even. The forging was carried out using a hydraulic press of type LYF-800SA. Before hot forging, the top and bottom die need to be heated, and the ceramic heating pad was placed in the middle of the two dies. The hydraulic press was controlled so that the top and bottom die just contact the heating pad, then the heating pad was turned on and the insulation

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cotton was wrapped around the die with insulation cotton. When the die temperature raised above 300 °C, the insulation cotton and heating pad were removed. To protect the die and smooth filling and pressurization process and reduce the ejection force, the die surface was sprayed oil-based graphite. The forged trackboards were heat treated. First, the material was placed in a box resistance furnace for solution treatment. Then, the material was immediately removed by warm water quenching, and then placed in a drying furnace for aging treatment. After a period of time, air cooling was taken out. After the heat treatment, the tracked plates were cut at the middle and edge of the test pieces, followed by coarse grinding, fine grinding, rough polishing, and fine polishing. The metallographic structure was observed using a Zeiss Axiovert 200 MAT optical microscope. The track shoe was cut at the cylinder position and the tensile pattern was processed. The tensile properties were evaluated using a GTM-9100 microcomputer-controlled universal material testing machine. JSM-6510 scanning electron microscope was used to observe fracture morphology. The hardness of the sample was measured using an HB-3000 Electronic Brinell Hardness Tester.

3 Results and Discussion 3.1 Formability Figure 3 shows the effects of various slug sizes and forging pressure on the formability of the trackboard by numerical simulation. As can be seen from Fig. 3, the magnitude of the pressure and the size of the slug during the forging greatly affect the final formability. Comparing Fig. 3a, b, when the pressure reaches 400 t, the calculation is completed. At this time, the die cavity has not been fully filled yet, and the metal slug at the edge cannot flow upward at this pressure. When the pressure reaches 800 t, the slug has filled the entire die cavity with good formability, indicating that the greater the pressure, the better the formability.

(a) No.1

(b) No.2

(c) No.3

Fig. 3 Numerical simulation results of trackboard formability with various slug sizes and forging pressures: a slug size 120 × 120 × 41, pressure 400 t, b slug size 120 × 120 × 41, pressure 800 t, c slug size 110 × 110 × 50, pressure 800 t

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Fig. 4 Three-pass forged SiC/2014 composite trackboard with different SiC contents: [2.5 wt% SiCp (left) and 10 wt% SiCp (right)]

Comparing Fig. 3b, c, even if the pressure reaches 800 t, No. 3 slug edge cannot be fully filled, and mainly because the slug is too thick. But if the deformation is too large, not only the flow stress is too large, but also the ductility of the material is insufficient. Cracks take place in areas where stress concentration occurs. Therefore, in the experimental process, the slug size was selected 120 mm × 120 mm × 41 mm, and the forging pressure was as close as possible to 800 t. In the practical experiment process, composite materials with SiC particles content of 2.5 and 10 wt% were coded and placed in a box type resistance furnace for heating. The temperature was raised to 480 °C for 1 h and the temperature was maintained for 2 h. When the die was heated to above 300 °C, hot forging began. With reference to the numerical simulation of the process parameters, the forging pressure of the composite material was gradually increased from 500 to 600 t, and the die cavity can be filled. The slug itself may be thicker than expected, and the forging left more flashes after a single forging. The thickness of the part center was required to be 5 mm. After the measurement, it was found that the thickness of the primary forging was more than 8 mm. Therefore, the flash was cut off and reheated for multiple forgings. Finally, forgings with better formability were obtained. Figure 4 shows the three-pass SiC/2A14 composite forgings containing 2.5 wt% SiCp (left) and 10 wt% SiCp (right), respectively.

3.2 Microstructure and Mechanical Properties The forgings shown in Fig. 4 were cut from the center, respectively, and separately numbered as 2.5-1, 2.5-2, 10-1, and 10-2, and then heat treated, as shown in Table 2. Samples were taken from the middle cylinder and the tensile specimens were machined to test their mechanical properties. Samples were cut and polished, and the microstructures of the samples were observed, as shown in Fig. 5. From the microstructure and mechanical properties after heat treatment, the performance of the forging composite trackboard has reached the strength and elongation

46 Table 2 Treat treatment process

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Number

Solid solution system

Aging system

2.5-1

Heat up to 500 °C for 1 h, hold for 5 h, warm water quench

Heat up to 165 °C, hold for 8 h, air cooling

2.5-2

Heat up to 500 °C for 1 h, hold for 5 h, warm water quench

Heat up to 165 °C, hold for 12 h, air cooling

10-1

Heat up to 500 °C for 1 h, hold for 5 h, warm water quench

Heat up to 165 °C, hold for 8 h, air cooling

10-2

Heat up to 500 °C for 1 h, hold for 5 h, warm water quench

Heat up to 165 °C, hold for 12 h, air cooling

Fig. 5 Optical microstructure of the SiC/2A14 composite forged trackboard: a 2.5-1, b 2.5-2, c 10-1, and d 10-2

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Fig. 6 Hardness of the SiC/2A14 composite forged trackboard

requirements (tensile strength ≥400 MPa, elongation >4%). It is clearly shown in Fig. 5 that the distribution of SiCp in the matrix is relatively uniform, the agglomeration phenomenon is not obvious, the particles and the matrix are combined well, and the composite material has no obvious organizational defects. The mechanical properties are shown in Figs. 6 and 7. The heat treatment aging time of 8 and 12 h has some effects on the mechanical properties, the tensile strength and hardness slightly increase, and the elongation decreases slightly. Comparing the mechanical properties of 2.5 wt% SiCp/2A14 and 10 wt% SiCp/2A14 composite samples, it can be found that the hardness and tensile strength of the latter are higher and the elongation does not change significantly. The appropriate SiCp content is 10 wt%, and the final heat treatment process is to heat up to 500 °C for 1 h, hold for 5 h, warm water quenching, and then to heat up to 165 °C, hold for 8 h, air cooling. Figure 8 shows the tensile fracture surface of the SiC/2A14 composite forged trackboard. It is noted that dimples with sharper outlines and smaller dimensions are distributed on the alloy substrate. However, there are some clean and flat platforms in the middle, and there is basically no matrix alloy organization, and it can be inferred that the large-particle SiC has undergone sensible cracking. The large-size SiC particles have many crystal defects, and the interface area between the matrix and the particles is large, and the load transmitted through the interface is large, leading to a greater understanding tendency. In the large deformation process of forging, it is difficult for large-size particles to cooperate with the flow of the matrix alloy, and stress concentration tends to occur at the particle aggregation to cause cracking of the particles [6]. The composite material is a hybrid type fracture, and its ductility and strength can meet the user requirements.

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Fig. 7 Tensile strength and elongation of the SiC/2A14 composite forged trackboard

(a)

(b)

(c)

(d)

Fig. 8 Dimples (red circles) and flat platforms (blue rectangles) tensile fracture of the SiC/2A14 composite forged trackboard: a 2.5-1, b 2.5-2, c 10-1, and d 10-2

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4 Conclusions (1) The simulated results showed that the size and quality of the blank had a great influence on the forming. Under the same blank volume, the thickness of the blank was too large, resulting in too much flash forgings, and the edge of the track shoe was not full and cracked. Formability of 800 t forming pressure at the same forming speed was better than 400 t, and the simulated results were in good agreement with the test results. (2) The hardness and tensile strength of SiCp/2A14 composite track shoe increased with the increase of silicon carbide content and aging time, while the elongation decreased with the increase of aging time. (3) For the SiCp/2A14 composite trackboard, the optimal hot forging process parameters were: slug size 120 mm × 120 mm × 41 mm, pressing speed 25 mm/s, die temperature 300 °C, and slug temperature 480 °C; Considering the stability of mechanical properties, the track shoe was suitable for selecting 10 wt% SiCp content. The heat treatment process is as follows: the solid solution system was heated from 1 h to 500 °C, kept for 5 h, quenched by warm water, and then heated to 165 °C for 8 h, air cooled.

References 1. M.A. Xavior, J.P.A. Kumar, Machinability of hybrid metal matrix composite—a review. Procedia Eng. 174, 1110–1118 (2017) 2. T.S. Srivatsan, I.A. Ibrahim, F.A. Mohamed et al., Processing techniques for particulatereinforced metal aluminium matrix composites. J. Mater. Sci. 26(22), 5965–5978 (1991) 3. G.H. Wu, Development challenge and opportunity of metal matrix composites. Acta Materiae Compositae Sin. 31(5), 1228–1237 (2014) 4. J.A. Liu, S.S. Xie, Application and Technology Development of Aluminum Alloy Materials (Metallurgical Industry Press, Beijing, 2004) 5. Z.L. Zhang, Z.F. Zhang, J. Xu, H. Zhang, W.M. Mao, Effect of mechanical stirring and air pressure on the fluidity of SiCp/A357 composites. Mater. Sci. Forum 898, 1000–1006 (2017) 6. B.L. Xiao, J. Bi, M.J. Zhao et al., Effect of SiCp size on tensile properties and fracture mechanism of aluminum matrix composites. Acta Metall. Sin. 38(9), 1006–1008 (2002)

Effect of Boron Addition Methods on Microstructure and Mechanical Properties of a Near-α Titanium Alloy Yingying Liu, Lihua Chai, Xiaozhao Ma, Yapeng Cui, Ziyong Chen and Zhilei Xiang

Abstract This work investigated the effect of boron addition methods on microstructure and mechanical properties of a near-α titanium alloy. Ti–6.5Al–2.5Sn–9Zr–0.5Mo–1W–1Nb–0.25Si was used as the matrix, and 0.3 wt% TiB2 and 0.1 wt% B were added, respectively. The results show that the addition of trace boron forms TiB whiskers on the prior β grain boundaries and leads to significant refinement of the microstructure in the based alloy. And, the refining effect of the 0.3 wt% TiB2 and 0.1 wt% B on the base alloy is similar. At room temperature, the strength of the boron-containing alloys has a certain increase, but the elongation drops slightly. Through study on the microstructure of tensile strained specimens, it was found that the increase of tensile strength of the boron-containing alloys is the combination of the base microstructure and the whisker bearing, while the ductility drops significantly is mainly attributed to the cracking of TiB phase.

1 Introduction Titanium and titanium alloys which hold the advantages of high specific strength, favorable corrosion resistance and low-temperature performance, high thermal strength, etc., have become a kind of critical structural materials in the aerospace industry, and moreover, have displayed considerable application potential for aeroengine heat-enduring parts owing to superior high-temperature performance compared with aluminum alloys and magnesium alloys [1–3]. At present, the main factor limiting the development of high-temperature titanium alloys to higher temperatures is the matching of thermal strength and thermal stability of the alloys [4]. In order to further improve the use temperature of the high-temperature titanium alloys, a large number of materials researchers have turned to research on composite materials, Whisker-reinforced titanium-based composites are typical representatives. However, the excessive reinforcement phase in the composite will increase the strength and Y. Liu · L. Chai (B) · X. Ma · Y. Cui · Z. Chen · Z. Xiang College of Materials Science and Engineering, Beijing University of Technology, Beijing 100124, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_6

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lead to an excessive decrease of the matrix plasticity. Therefore, in order to comprehensively match the strength and plasticity of the alloy, boron-containing alloys have attracted interest in recent years. Compared to traditional alloys and composites, boron-containing alloys ensure high strength and excellent wear resistance without excessively reducing alloy plasticity, and the overall performance is well matched [5–9]. Based on the traditional Ti–Al–Sn–Zr–Mo–Si near-α type high-temperature titanium alloy, this paper designed a high-temperature titanium alloy. Study on the effect of trace boron on solidification microstructure and properties of high-temperature titanium alloy by adding different weight fractions of TiB2 and boron powder. It is expected that the performance of high-temperature titanium alloys will be further enhanced through the addition of trace amounts of boron to meet the engineering application requirements. The present work was aimed to study the effect of boron addition methods on as-cast microstructure and mechanical properties of a hightemperature titanium alloy.

2 Materials and Experimental Near-α titanium alloy with a nominal composition of Ti–6.5Al–2.5Sn–9Zr0.5Mo–1W–1Nb–0.25Si alloy (wt%) as initial alloy components. The raw materials to prepare this near-α titanium alloy with and without boron additions were sponge titanium (grade 0), aluminum thread (99.5%), sponge zirconium (99.5%), silicon powder (99.99%), boron powder (99.9%), and Al–Mo, Ti–Sn, Al–Nb, Al–W, and Al–TiB2 master alloys. Table 1 lists the nominal composition of alloys under study. 4 kg bar of the boron-containing titanium with an approximate size of ∅60 × 150] mm was melted in a laboratory vacuum induction melting furnace under argon atmosphere. The ingots were remelted three times to ensure good homogeneity. The matrix alloy was obtained by remelting of the initial alloy to produce the ascast condition in the same manner. Slices were cut from the center of the ingot for microstructure examination. The Ti–6.5Al–2.5Sn–9Zr–0.5Mo–1W–1Nb–0.25Si and Ti–6.5Al–2.5Sn–9Zr–0.5Mo–1W–1Nb–0.25Si–(0.1B or 0.3TiB2 ) alloys are designated as TA6.5, TA6.5–0.1B, and TA6.5/0.3TiB2 in this work. A FEI Quanta 650 scanning electron microscopy (SEM) or imagerA2m ZEISS optical microscopy (OM) was used for the examination of microstructures and fracture surfaces of the failed specimens. Prior to OM and SEM studying, the specimen surfaces were polished and etched. The etchant composition is 5%HNO3 + 3%HF + 92%H2 O. The size of prior β grains, the α/β colony size was determined by linear intercept method. For tensile tests, the flat specimens were cut from the ingots (Fig. 1). Three specimens were tested in each condition, and the strains were measured using a 25 mm gauge length extensometer. The tensile tests at room temperature were carried out on an Instron machine with initial strain rate of 6 × 10−4 s−1 and the ultimate

Bal Bal Bal

Ti–6.5Al–2.5Sn–9Zr–0.5Mo–1W–1Nb–0.25Si/0.3TiB2

Ti–6.5Al–2.5Sn–9Zr–0.5Mo–1W–1Nb–0.25Si–0.1B

Ti

6.5

6.5

6.5

Al

Composition, wt%

Ti–6.5Al–2.5Sn–9Zr–0.5Mo–1W–1Nb–0.25Si

Alloy

Table 1 Chemical compositions of the alloys under study

2.5

2.5

2.5

Sn

9.0

9.0

9.0

Zr

0.5

0.5

0.5

Mo

1.0

1.0

1.0

Nb

0.25

0.25

0.25

Si

1.0

1.0

1.0

W

0.1

0.1

–

B

Effect of Boron Addition Methods on Microstructure … 53

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Fig. 1 Schematic tensile sample at room temperature

tensile strength, σUTS , the yield strength, σ0.2 , and elongation to rupture, δ, were determined from the tests. All specimens for microstructural examination and mechanical testing were prepared by electrospark cutting by fine grinding of work surfaces.

3 Results and Discussion 3.1 Effect of Boron Addition on Microstructure Figure 2a represents BSE images of the base alloy in as-cast conditions. Coarse prior β-grains and α/β colonies are observed. The prior β grain boundaries are complete and there is a continuous grain boundary α phase distributes in it. The α colonies in β-grains are arranged in a parallel layer with the same orientation and the needles are distributed in different orientations. The 0.3 wt% TiB2 addition results in the formation of uniformly distributed TiB whiskers with random orientation, which refines the matrix microstructure (Fig. 2c). Figure 3 shows high-magnification secondary electron images of boroncontaining titanium alloys in as-cast conditions. It can be seen that TiB whiskers distribute at the prior β grain boundary (Fig. 3a). The presence of borides leads to distortions of the prior β grain boundaries while limiting the growth of α/β colonies in β-grains, so the size of the α/β colonies becomes shorter and the α lath width exhibit relatively finer. One can see that the additions of 0.3 wt% TiB2 and 0.1 wt% B to the matrix have similar effects on the refinement of prior β-grains (Fig. 2b, c). The changes of grain size and the α/β colonies size are basically similar, and there is no obvious difference between the two microstructures. Image analysis software was used to quantitatively evaluate the prior β-grains size d and the thickness of α lath of boron-containing titanium alloys with two different boron additions and the matrix. The test results are shown in Fig. 4. One can see

Effect of Boron Addition Methods on Microstructure …

55

Fig. 2 Effects of TiB2 /B on microstructure of as-cast 5 a TA6.5, b TA6.5/0.3TiB2 , and c TA6.5–0.1B alloys

Fig. 3 SEM micrographs of TA6.5/0.3TiB2 : a lamellar structure, b magnification of square part in a

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Fig. 4 Prior β-grains size and α lath thickness of different boron additions

that the alloys containing 0.3 wt% TiB2 and 0.1 wt% boron show finer β-grains with prior β-grain size (451.45 ± 105.2 μm) reaching up to 156.69 ± 26.3 μm and 144.70 ± 32.8 μm and the thickness λ of α lath (4.20 ± 0.95 μm) reaching up to 1.33 ± 0.64 μm and 1.70 ± 0.37 μm, respectively. From which, it can be concluded that the refining effect of TiB2 and element boron on the prior β-grains of the matrix alloy is similar.

3.2 Reaction of B/TiB2 in Titanium Alloys In the titanium alloy system, when the content of boron is small, TiB2 cannot exist stably, and reacts with Ti to form TiB which is also a reinforcing phase [10]. Therefore, TiB2 is usually used as a source for TiB generation. As can be seen from the Ti–B binary phase diagram [11], TiB can be generated in situ by the following two reactions: Ti + B → TiB TiB2 + Ti → 2TiB In order to confirm that TiB2 and boron have been transformed into TiB in the titanium matrix, XRD analysis of 0.3 wt% TiB2 and 0.1 wt% B were tested, and the results are shown in Fig. 5. It can be seen that the two alloys are mainly composed of Ti and TiB and no metastable boride was found.

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Fig. 5 X-ray diffraction patterns of the boron-containing titanium alloys

3.3 Grain Refinement Mechanism 3.3.1

Refinement Mechanism of the Prior β Grains

Numerous researches reported [12–14] that the addition of trace boron in titanium alloy can significantly refine the grain sizes. Zhu pointed out that the boron added to titanium alloy forms TiB, and believed that the TiB acts as a heterogeneous nucleation point refines the prior β grains [15]. As can be seen from the Ti–B binary phase diagram, when the boron content is lower than the eutectic composition, the prior β phase first precipitates when the temperature is lower than the liquidus during solidification, and then the temperature continues to drop to 1540 °C, an eutectic reaction occurs to generate TiB. Therefore, Tamirisakandala [16] negated the possibility of TiB as a heterogeneous nucleus to refine the grains. For the as-cast alloy, grain refinement can generally be caused by undercooling, undercooling of components, or introduction of heterogeneous nucleation sites. Tamirisakandala [16] proposed a mechanism that supercooling of the boron-containing titanium alloys arouses grain refinement, as shown in Fig. 6. When the temperature is higher than the liquidus during solidification, the boron atoms are dissolved in the liquid phase. Due to the low solubility of boron in beta phase (0.02 wt%), the prior β grains begin to nucleate in the melt when the temperature drops below the liquidus, therefore, part of the solute boron at the nucleated β grains will be pushed to the front of the solid–liquid interface, causing constitutional supercooling to make the solid–liquid interface unstable. The new phase in the liquid continuously nucleates. When the temperature continues to fall to the eutectic temperature, the boron atoms accumulated at the front of the β phase in the melt combine with the remaining liquid and eventually forms TiB uniformly distributed on the β grain boundary.

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Fig. 6 Grain refinement mechanism in the boron-containing titanium alloys

3.3.2

Refinement Mechanism of α Lath

Slow decrease of the temperature in the solidification process will result in the transformation of β → α allotrope. The α phase precipitates from the β phase according to the {011}β//{0001}α and [111]//[11–20] α Burgers orientation relations [17]. The fine grain boundary generated by the prior β grain provides more heterogeneous nucleation sites for the nucleation of α phase, so that the nucleation rate of α phase increases, and the α lath is refined. In addition, TiB is also present in the melt, it is not excluded that TiB as a heterogeneous nucleation site refines the α lath. And due to the refined of the prior β grains, the α lath growth is also limited to a certain extent. When TiB2 is used as a boron source to introduce a titanium alloy matrix, it is known that a small amount of TiB2 is unstable in the titanium alloy matrix. During the solidification process, TiB2 has been transformed to TiB at a temperature higher than the liquidus. Therefore, grain refinement of the alloy mainly depends on the existence state of TiB in the melt [18]. At higher temperatures (about 1700 °C), TiB is in a thermodynamically unstable state in the melt, cracking produces boron atoms and Ti radicals, and boron atoms are pushed to the solid–liquid interface initiating constitutional supercooling to refine the original β-grains during the solidification. With the temperature dropping, TiB in the melt gradually changes to a stable state and no cracking. Therefore, TiB can be used as a heterogeneous nucleus to refine the prior β grains.

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3.4 Mechanical Properties 3.4.1

Tensile Properties at Room Temperature

Results of the tensile properties of the as-cast alloys with two different boron addition methods at room temperature are summarized in Table 2. The tensile curves of the base alloy and boron-containing alloys are shown in Fig. 7. As can be seen from it, the boron-containing alloys show a higher strength as compared to the base alloy. The tensile strength and yield strength at room temperature increase by 41.2 and 48.8 MPa, respectively, after adding 0.3 wt% TiB2 to the base alloy, and the elongation decreases to 5.6%. After adding 0.1 wt% B to the base alloy, the room temperature tensile strength and yield strength of the alloy increase by 63.8 and 64.3 MPa, respectively, and the elongation decreases to 5.9%. Comparing two boron addition methods, the addition of TiB2 and element boron has increased the strength of the as-cast alloys, but the increase is not significant and there is slight decrease in ductility in both the boron-containing alloys. Similarly, Chandravanshi [19] obtained a similar conclusion in the study of the room temperature tensile properties of Ti–1100B boron-containing titanium alloys. Therefore, from the perspective of improving the mechanical properties of the as-cast matrix alloys, the additional effects of TiB2 and element boron are basically similar. Comparing the properties of the boron-containing alloys in other literatures, it can be seen that the performance of the TA6.5–0.1B alloy is significantly better than other boron-containing alloys, and the strength of the alloy is improved under the premise of ensuring plasticity.

3.4.2

Fracture Behavior

Figure 8 illustrates the fracture surfaces of tensile specimens of the base alloy and the boron-containing alloys with different boron addition methods after tensile straining at room temperature. There are some dimples distributed in the fracture surfaces of the base alloy, indicating that which was failed by microvoid coalescence and has higher plastic at room temperature (Fig. 8a). While 0.3 wt% TiB2 was added to the

Table 2 Tensile properties of the boron-containing alloys in as-cast condition

Alloy type

Tensile properties σUTS /MPa

σ0.2 /MPa

δ/%

TA6.5

1019.1

885.4

8.7

TA6.5/0.3TiB2

1060.3

934.2

5.6

TA6.5–0.1B

1082.9

949.7

5.9

Ti1100–0.2B [19]

1037.0

938.0

1.3

Timetal 685–0.2B [20]

917.0

828.0

2.5

T64–0.1B [21]

965.8

900.0

7.3

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Fig. 7 Tensile curves of the boron-containing alloys with different boron additions in the as-cast condition at room temperature

matrix, the fracture surfaces show very shallow dimples and mainly consist of TiB cleavage plane and matrix torn edges, belonging to quasi-cleavage fracture, which shows lower plastic as compared to the matrix alloy (Fig. 8b). The fracture surfaces of the alloy with boron content of 0.1 wt% are similar to the fracture structure with 0.3 wt% TiB2 alloy, is mainly dominated by TiB cleavage surface and matrix torn edges. Plasticity closes to the alloy with 0.3 wt% TiB2 (Fig. 8c). The behavior of the boron-containing alloys at room temperature is attributed to the random distribution of the whiskers as well as the brittle of the near-α titanium alloy matrix in the as-cast condition [22]. The whiskers undergo shear stress along the matrix/whiskers interface during the tensile test, while the strength of the boride whiskers is affected by their orientation, morphology (aspect ratio) and volume fraction in the matrix, the random distribution of whiskers in the as-cast condition has low strength and thus causes to premature fracture initiation cracks [23]. With the loading stress increases, these cracks continue to expand, eventually leading to matrix fracture. The increase in plasticity caused by the microstructural refinement resulting from the supercooling of element boron in as-cast condition is eventually offset by the embrittlement of the whiskers. Figure 9 illustrates the flat surfaces of tensile strained specimens of the 0.1 wt% B alloy near the fracture zone. The TiB whiskers, which are nearly parallel to the tensile direction, were broken into several pieces, and pores appeared on the TiB particle–matrix interface. However, no debonding between the reinforcing phase and the matrix was observed, indicating that high adhesion strength of coherent boundaries exists between Ti-matrix and TiB whiskers (Fig. 9a). The TiB fracture phenomenon was not observed in the direction perpendicular to the tensile axis, and the crack propagated in the transverse direction. This is due to the fact that whiskerlike TiB bonds between the atoms in [010] direction are weaker than [001] direction [24]. The stress field appears between the reinforcement phase and the matrix due to shearing during the tensile test. As the plasticity of the matrix is higher than that of the reinforcement phase, the stress concentration continuously accumulates in the tensile test, cracks occur above the critical stress, and then the crack propagates in the horizontal direction.

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Fig. 8 Fracture surfaces of specimens after tensile at room temperature in as-cast condition: a TA6.5, b TA6.5/0.3TiB2 , and c TA6.5–0.1B

When the content of boron added in the matrix is minor, the increase of the strength of the boron-containing titanium alloys is mainly due to the effect of refinement of the prior β grains and the α lath caused by t element-induced boron undercooling, The decrease in plasticity is mainly due to the cracking of the reinforcing phase TiB introduced by micro-boron during plastic deformation [25]. With the increase of boron content, the degree of refinement gradually increases. When the boron content is 0.1 wt%, boron-containing titanium alloy has the best overall performance. The TiB whisker generated in the in situ reaction has a good interface with the matrix, which can transfer the applied load of the matrix, thus increases the strength of the matrix alloy.

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Fig. 9 Morphologies of deformation zone around tensile fracture at room temperature of TA6.5–0.1B alloy

4 Conclusions (1) The addition of trace boron in near-α Titanium alloy can cause significant refinement of the prior β-grains and the α-lath, and the 0.3 wt% TiB2 and 0.1 wt% B had similar refinement effects. (2) The tensile strength had a certain degree of improvement after adding 0.3 wt% TiB2 and 0.1 wt% B compared with the matrix alloy, while the elongation was decreased slightly. (3) The fracture behavior showed that when the small amount of boron was added, the increase of the strength of the boron-containing alloys was mainly attributed to the refinement of the matrix structure caused by elemental boron.

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References 1. M. Li, J. Luo, et al., Precision Forging of Titanium Alloy. Science Press (2016) 2. M.R. Winstone, A. Partridge, J.W. Brooks, The contribution of advanced high-temperature materials to future aero-engines. Proc. Inst. Mech. Eng. Part L J. Mater. Des. Appl. 215(2), 63 (2001) 3. J. Cai, C. Cao, Alloy design and application expectation of a new generation 600 °C high temperature titanium alloy. J. Aeronaut. Mater. 34(4), 27 (in Chinese) 4. J. Cai, Z. Li, J. Ma, Research and development of 600 °C high temperature titanium alloys for aeroengine. Mater. Rev. (2005) 5. C.J. Boehlert, C.J. Cowen, S. Tamirisakandala et al., In situ scanning electron microscopy observations of tensile deformation in a boron-modified Ti–6Al–4V alloy. Scripta Mater. 55(5), 465–468 (2006) 6. Y. Qin, L. Geng, D. Ni, Dry sliding wear behavior of extruded titanium matrix composite reinforced by in situ TiB whisker and TiC particle. J. Mater. Sci. 46(14), 4980–4985 (2011) 7. Z. Zhang, J. Qin, Z. Zhang et al., Microstructure effect on mechanical properties of in situ, synthesized titanium matrix composites reinforced with TiB and La2 O3 . Mater. Lett. 64(3), 361–363 (2010) 8. B.J. Choi, I.Y. Kim, Y.Z. Lee et al., Microstructure and friction/wear behavior of (TiB + TiC) particulate-reinforced titanium matrix composites. Wear 318(1–2), 68–77 (2014) 9. J.S. Kim, K.M. Lee, D.H. Cho et al., Fretting wear characteristics of titanium matrix composites reinforced by titanium boride and titanium carbide particulates. Wear 301(1–2), 562–568 (2013) 10. S. Gorsse, Y.L. Petitcorps, S. Matar et al., Investigation of the Young’s modulus of TiB needles in situ produced in titanium matrix composite. Mater. Sci. Eng., A 340(1–2), 80–87 (2003) 11. J.L. Murray, P.K. Liao, K.E. Spear, The B − Ti (Boron-Titanium) system. Bull Alloy Phase Diagr 7(6), 550–555 (1986) 12. G. Zorn, Glass formation in boron-containing alloys by mechanical alloying*: Zeitschrift für Physikalische Chemie. Zeitschrift Für Physikalische Chemie, 157(Part_1), 203–208 (1988) 13. J.B. Jergenson. Preparation of liquid metal source structures for use in ion beam evaporation of boron-containing alloys (1986) 14. R. Sarkar, P. Ghosal, K. Muraleedharan et al., Effect of boron and carbon addition on microstructure and mechanical properties of Ti-15-3 alloy. Mater. Sci. Eng., A 528(13–14), 4819–4829 (2011) 15. J. Zhu, A. Kamiya, T. Yamada et al., Influence of boron addition on microstructure and mechanical properties of dental cast titanium alloys. Mater. Sci. Eng. A 339(1–2), 53–62 (2003) 16. S. Tamirisakandala, R.B. Bhat, J.S. Tiley et al., Grain refinement of cast titanium alloys via trace boron addition. Scripta Mater. 53(12), 1421–1426 (2005) 17. W.G. Burgers, On the process of transition of the cubic-body-centered modification into the hexagonal-close-packed modification of zirconium. Physica 1(7), 561–586 (1934) 18. T.T. Cheng, The mechanism of grain refinement in TiAl alloy by boron addition—an alternative hypothesis. Intermetallics 8(1), 29–37 (2000) 19. V.K. Chandravanshi, R. Sarkar, S.V. Kamat et al., Effect of boron on microstructure and mechanical properties of thermomechanically processed near alpha titanium alloy Ti-1100. J. Alloy. Compd. 509(18), 5506–5514 (2011) 20. V.K. Chandravanshi, R. Sarkar, P. Ghosal et al., Effect of Minor Additions of Boron on Microstructure and Mechanical Properties of As-Cast Near α, Titanium Alloy[J]. Metall Mater Trans A 41(4), 936–946 (2010) 21. I. Sen, S. Tamirisakandala, D.B. Miracle et al., Microstructural effects on the mechanical behavior of B-modified Ti–6Al–4V alloys. Acta Mater. 55(15), 4983–4993 (2007) 22. S. Gorsse, D.B. Miracle, Mechanical properties of Ti-6Al-4V/TiB composites with randomly oriented and aligned TiB reinforcements. Acta Mater. 51(9), 2427–2442 (2003) 23. W. Lu, Z. Di, X. Zhang et al., Microstructural characterization of TiB in in situ synthesized titanium matrix composites prepared by common casting technique. J. Alloy. Compd. 327(1), 248–252 (2001)

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24. F. Ma, T. Wang, P. Liu et al., Mechanical properties and strengthening effects of in situ, (TiB + TiC)/Ti-1100 composite at elevated temperatures. Mater. Sci. Eng. A 654, 352–358 (2016) 25. V.M. Imayev, R.A. Gaisin, E.R. Gaisina, et al. Microstructure, processing and mechanical properties of a titanium alloy Ti-20Zr-6.5Al-3.3Mo-0.3Si-0.1B. Mater. Sci. Eng. A (2017)

Effect of Ultrasonic Treatment on Microstructure and Properties of Aluminum Alloy Rod Prepared by Continuous Casting and Rolling Nan Zhou, Shuncheng Wang, Zhimin Zhang, Yuanhui Zhai and Guangyao Zhong Abstract The Al–0.78Mg–0.69Si–0.11Ce–0.06La aluminum alloy rod was prepared by ultrasonic-assisted continuous casting and rolling process, and the effects of ultrasonic treatment on the microstructure, electrical conductivity, and tensile properties of the continuous casting billet and the continuous casting and rolling rod were investigated. The results showed that ultrasonic treatment could purify, homogenize, and refine the aluminum alloy liquid, eliminate the coarse dendrite of continuous casting billet, restrain the element segregation, and significantly improve the strength, plasticity, and conductivity of the aluminum alloy rod. When the ultrasonic frequency was 20 kHz and the power was 300 W, the tensile strength, elongation, and electrical conductivity of the aluminum alloy rod were 214.5 MPa, 9.6%, and 55.9% IACS, respectively. Compared with the continuous casting and rolling aluminum alloy rod without ultrasonic assistance, the tensile strength, elongation, and conductivity were increased by 13.6, 41.2, and 2.4%, respectively.

1 Introduction With the rapid development of the economy and the continuous improvement of the living standard, the demand for electric power is increasing rapidly, so the transmission line is increasingly developing toward the direction of large capacity, which requires the increase of transmission capacity of conductors [1]. At present, the transmission wire in China is mainly the traditional steel core aluminum strand, which is relatively high in strength, but large in electric energy loss, low in heat resistance, and greatly limited in transmission capacity [2]. In order to reduce the power loss of transmission lines and improve the efficiency of power utilization, the State Grid Corporation of China has vigorously promoted the application of high-strength and N. Zhou (B) · S. Wang Gaungdong Institute of Materials and Processing, Guangzhou 510650, China e-mail: [email protected] Z. Zhang · Y. Zhai · G. Zhong Guangdong Shine Cables Co., Ltd., Qingyuan 511500, China © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_7

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high-conductivity, high-strength, and heat-resistant aluminum alloy conductors in recent years [3, 4]. Al–Mg–Si series aluminum alloy is a kind of deformed aluminum alloy which can be strengthened by heat treatment. It has the advantages of high strength and good plasticity, which is suitable for manufacturing aluminum alloy conductor with high strength and high conductivity for long-span long-distance transmission lines. However, due to the high content of Mg and Si in Al–Mg–Si system, the coarse dendrite is easily formed in the continuous casting billet, which leads to the formation of element segregation and pore, and furthermore reduces the plasticity of the continuous casting billet, which would cause the fracture of the continuous casting billet in the rolling process and result in the reduction of the production efficiency and the rate of final products [5]. Therefore, in this paper, the Al–0.78Mg–0.69Si–0.11Ce–0.06La aluminum alloy rod was fabricated by ultrasonic-assisted continuous casting and rolling process. The effects of ultrasonic treatment on the microstructure, conductivity and tensile properties of the continuous casting billet and continuous casting and rolling rod were studied.

2 Experimental Al–Mg–Si–Ce–La system aluminum alloy was prepared by industrial pure aluminum ingot, pure magnesium ingot, crystalline silicon, and lanthanum cerium mixed aluminum master alloy AlRE10, with the chemical composition (wt%) of 3.5 lanthanum, 6.5 cerium, and balanced aluminum. The chemical composition of Al–Mg–Si–Ce–La alloy which was determined by ARL460 photoelectric direct reading spectrometer is shown in Table 1. The experimental equipment included a 15 ton aluminum melting furnace, a round tilting heat preservation furnace with permanent magnetic stirring, a wheel belt continuous caster imported from France and a 15 stand Y type three-roller continuous rolling mill. The diameter of the crystal wheel in the continuous casting machine was 1600 mm, the cross section area of the continuous casting ingot was 2400 mm2 , the linear speed of the continuous casting ingot was 0.2 m/s, the cooling water flow was 60 m3 /h, and the finishing rolling speed of the continuous rolling mill was 6.2 m/s. The alloy was melted in the aluminum melting furnace. First, the alloy liquid was stirred by a permanent magnetic stirring which can uniform the composition of the aluminum alloy liquid, and then treated by injection refining to remove the gas and slag in the alloy liquid. After static holding, the alloy liquid was poured into the continuous casting machine. With the assistance of the ultrasonic wave, the aluminum alloy rod with the diameter of 9.5 mm was prepared by continuous casting and rolling.

Table 1 The chemical composition of the experimental alloy (wt%)

Elements

Mg

Si

Ce

La

Fe

Al

Composition

0.78

0.69

0.11

0.06

0.09

Bal.

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The ultrasonic frequency was 20 kHz, the ultrasonic power was, respectively, 0, 100, 200, and 300 W, the casting temperature was 710 °C, the start rolling temperature was 490 °C, and the final rolling temperature was 320 °C. After grinding, polishing and etching, the microstructures were observed on an OLYMPUS-M10 optical microscope (OM), and the microstructure was observed by a PYM-150 environmental scanning electron microscope (SEM) equipped with energy dispersive spectrometer (EDS). The phase composition of the alloy was analyzed by a WDM X-ray diffractometer (XRD). The electrical conductivity was measured by a WCW-500 DC double arm bridge conductance instrument. The tensile strength and elongation were tested at room temperature on an NYP1000 electronic tensile tester, with the tensile speed at 2 mm/min.

3 Results and Discussion 3.1 Microstructure Figure 1 shows the microstructure of Al–0.78Mg–0.69Si–0.11Ce–0.06La alloy continuous casting billet after ultrasonic treatment at different power. It can be seen from Fig. 1 that without ultrasonic treatment, the microstructure of the continuous casting billet was composed of coarse dendrite, with the eutectic phase distributed between the dendrites, and the segregation of Mg and Si elements on the surface was serious, as shown in Fig. 1a. After ultrasonic treatment, the microstructure of the continuous casting billet refined, and with the increasing ultrasonic power, the grain morphology gradually changed from coarse dendrites to equiaxed grains, the size of which decreased gradually [6]. When the ultrasonic power increased to 300 W, the grain structure of the aluminum alloy continuous casting billet had totally changed into fine and uniform equiaxed crystal, as shown in Fig. 1d. The experimental results showed that the grain structure of aluminum alloy continuous casting billet could be significantly refined by ultrasonic treatment, and the microstructure of aluminum alloy continuous casting billet could refine from coarse dendrites to fine uniform equiaxed grains. Figure 2 shows the XRD pattern of Al–0.78Mg–0.69Si–0.11Ce–0.06La alloy continuous casting billet after ultrasonic treatment. Figure 3 shows the SEM image of the continuous casting billet. It can be seen from Fig. 2 that the second phases of the alloy continuous casting billet were mainly Mg2 Si phase and AlFeSiLaCe phase. It can be seen from Fig. 3 that bright contrasted granular compounds and short rod compounds can be observed in the SEM microstructure of the continuous casting billet. EDS results revealed that the bright contrasted granular compound was Mg2 Si phase, and the short rod compound was AlFeSiLaCe phase. Besides, it can also be seen from Fig. 3 that the Mg2 Si phase and AlFeSiLaCe phase were fine and uniformly distributed in the α-Al matrix.

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Fig. 1 Optical microstructures of the continuous casting billet with ultrasonic treatment at different power: a 0, b 100 W, c 200 W, and d 300 W Fig. 2 XRD pattern for the continuous casting billet of the aluminum alloy

Intensity, I/a.u.

α-Al Mg2Si AlFeSiLaCe

10

20

30

40

50

2θ / ( ° )

60

70

80

90

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Fig. 3 SEM image for the continuous casting billet of the aluminum alloy

Mg2Si

AlFeSiLaCe

3.2 Electrical Conductivity Figure 4 shows the electrical conductivity of Al–0.78Mg–0.69Si–0.11Ce–0.06La aluminum alloy continuous casting and continuous rolling rod treated with different ultrasonic power. It can be seen from Fig. 4 that the electrical conductivity of the continuous casting and rolling aluminum alloy circular rod was relatively low without ultrasonic treatment. With the adding of the ultrasonic treatment, the conductivity of the continuous casting and rolling increased significantly. The ultrasonic treatment had a deeply purification effect on the aluminum alloy liquid, which can eliminate the harmful effect as metal and nonmetallic inclusions hydrogen and gas hole on the conductivity of the aluminum alloy rod, and furtherly improve the conductivity of the aluminum alloy rod. As can also be seen from Fig. 4, with the gradual increase of the ultrasonic power, the conductivity of the aluminum alloy rod increased gradually. When the power was increased to 300 W, the conductivity of the aluminum alloy rod increased to 55.9% IACS. Compared with the aluminum alloy rod without ultrasonic treatment, the conductivity of the continuous casting and rolling aluminum alloy rod increased by 2.4%.

3.3 Mechanical Properties Figure 5 shows the tensile strength and elongation of the continuous casting and rolling Al–0.78Mg–0.69Si–0.11Ce–0.06La alloy rods with ultrasonic treatment at different power. As can be seen from Fig. 5, without ultrasonic treatment, the tensile strength and elongation of the aluminum alloy rod were lower because the microstructure of the aluminum alloy rod was coarse dendrite. In this condition, the tensile strength was 188.9 MPa and the elongation was 6.8%. With the addition of ultrasonic treatment, the microstructure of aluminum alloy continuous casting billet

Fig. 4 Electrical conductivity of the aluminum alloy rod with ultrasonic treatment at different power

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Electircal conductivity / %IACS

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Ultrasonic power / W

changed from coarse dendrites to fine uniform equiaxed grains, and ultrasonic treatment has a deep purification effect on aluminum alloy liquid. Therefore, the harmful effects of inclusions and pores on the tensile mechanical properties of the aluminum alloy rods can be eliminated efficiently, so that the tensile strength and elongation of the aluminum alloy rod can be improved simultaneously [7]. In addition, ultrasonic treatment can also inhibit the segregation of Mg and Si elements in aluminum alloy continuous casting billet, improve the homogeneity of composition distribution, refine the Mg2 Si phase and AlFeSiLaCe phase, and make the Mg2 Si phase and AlFeSiLaCe phase distribute uniformly in α-Al matrix, which could furtherly enhance the strength and ductility of the continuous casting and rolling aluminum alloy rod [8]. Besides, it can also be seen from Fig. 5 that the tensile strength and ductility of the aluminum alloy rod increased gradually with the increase of the ultrasonic power. When the ultrasonic power was 300 W, the tensile strength and elongation of the continuous casting and rolling aluminum alloy rod were 214.5 MPa and 9.6%, respectively. Compared with those without ultrasonic treatment, the tensile strength and elongation of the continuous casting and rolling aluminum alloy rod increased by 13.6 and 41.2%, respectively.

4 Conclusions (1) Ultrasonic treatment can purify and refine aluminum alloy liquid, eliminate the coarse dendrite and gas pore of aluminum alloy continuous casting billet, restrain element segregation, and improve the strength, ductility, and electrical conductivity of the continuous casting and rolling aluminum alloy rod. (2) When the ultrasonic frequency was 20 Hz and the power was 300 W, the tensile strength, elongation, and electrical conductivity of the continuous casting and

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Elongation / %

Tensile strength / MPa

Effect of Ultrasonic Treatment on Microstructure and Properties …

Tensile strength Elongation

Ultrasonic power / W Fig. 5 Tensile mechanical properties for aluminum alloy rod with different ultrasonic power

rolling aluminum alloy rod were 214.5 MPa, 9.6%, and 55.9% IACS, respectively. (3) Compared with the aluminum alloy rod without ultrasonic treatment, the tensile strength, elongation, and electrical conductivity of the continuous casting and rolling aluminum alloy rod increased by 13.6, 41.2, and 2.4%, respectively. Acknowledgements The authors would like to acknowledge the financial support of the Guangdong Academy of Sciences (2017GDASCX-0117, 2018GDASCX-0117), Guangdong Science and Technology Department (2017A070701029), and Qingyuan Municipal Bureau of Science and Technology (201601).

References 1. C. Huang, The application and development of aluminium and aluminium alloy for electrical purpose in cable field. Electr. Wire Cable 2, 10–15 (2013) 2. B. Liu, Q. Zheng, P. Dang et al., Development and applications of aluminium alloy in overhead lines. Electr. Wire Cable 4, 10–15 (2012) 3. Z. Li, C. Li, Y. Feng, Energy-saving effect analysis of energy-saving wire in transmission line. Electr. Power. Sci. Eng. 32(10), 28–33 (2016) 4. G. Ding, Z. Sun, Q. Zhang et al., Analysis on application of energy-saving conductors in transmission lines. Power Syst. Technol. 36(8), 24–30 (2012) 5. J. Hu, T. Zhou, Z. Li et al., Production status and development prospects of Al–Mg–Si alloy conductor. Light Alloy Fabr. Technol. 46(1), 5–8 (2018) 6. L. Zhang, X. Zhang, Effect of ultrasonic power on grain refinement in semi-continuous casting of 7050 aluminum alloy. Mater. Mech. Eng. 33(9), 54–57 (2009)

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7. H. Hu, K. Chen, L. Huang et al., Effect of ultrasonic melt treatment on microstructure and mechanical properties of Al-Zn-Mg-Cu alloy. Heat Treat. Metals 30(5), 43–46 (2005) 8. C. Shi, K. Shen, D. Mao, et al., Effects of ultrasonic treatment on microstructure and mechanical properties of 6016 aluminium alloy. Mater. Sci. Technol. https://doi.org/10.1080/02670836. 2018.1465514

Influence of High-Temperature Compression on Microstructure and Properties of Sintered Molybdenum Fu Wang, Zenglin Zhou, Yan Li, Zhilin Hui, Xueliang He and Xiaoying Fu

Abstract The high-temperature compression of sintered molybdenum was performed at 1200 °C and strain rate of 0.1 s−1 on Gleeble 3500 system, and the deformation amount reached to 50%. The density, Vickers hardness, microstructure, and texture of sintered and compressed molybdenum specimens have been measured and evaluated by Archimedes method, Vickers hardness tester, optical microscopy (OM), and X-ray diffraction (XRD), respectively. The results show that the relative density of Mo specimen increases from 95.9 to 99.4% after high-temperature compression, which is close to fully dense. The Vickers hardness HV10/10 significantly is raised from 174.5 ± 6.8 MPa to 224.1 ± 21.0 MPa. The percentage of hardness fluctuation increases from 7.8 to 18.7%, which indicates that the uniformity of hardness distribution is decreased, and the hardness distribution shows a certain regularity. Combined with optical microstructure, it can be found that there are hard-deforming zone (zone I), free-deforming zone (zone III), and easy-deforming zone (zone II) in the compressed specimen. The texture analysis shows that there is a certain intensity of 113//ND fiber texture in the sintered Mo specimen, while there are multiplex textures in the compressed one, including the {112}111, {001}100, {001}110, and the residual {113}111 sintering texture.

1 Introduction The refractory metal molybdenum has been widely used in manufacturing hightemperature components which are applied in high-end equipment manufacturing, petrochemicals, aerospace, nuclear industries, and other fields, due to its high melting point, high Young’s modulus, high creep resistance, low thermal expansion coefficient, excellent thermal and electrical conductivities, superior high-temperature strength and good resistance to thermal shock, as well as outstanding resistance to acid and alkali corrosion [1–5]. However, molybdenum, as a typical body-centered-cubic F. Wang · Z. Zhou (B) · Y. Li · Z. Hui · X. He · X. Fu Powder Metallurgy and Special Materials Research Department, GRIMAT Engineering Institute Co., Ltd, Beijing, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_8

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(BCC) metal, has ductile-to-brittle transition temperature (DBTT) which is close to room temperature, and the DBTT is very sensitive to interstitial impurity atoms (C, O, etc.) and the forming methods. Equiaxed crystalline molybdenum metal shows severe brittleness at room temperature [6, 7]. In addition, metal molybdenum and molybdenum alloys are easily oxidized, especially destructive oxidation occurs if the temperature exceeds 700 °C [8]. Mo also exhibits large resistance to deformation at low temperatures [9]. All these characteristics seriously restrict the development and application of molybdenum products. In industrial practice, refractory metal molybdenum products are mostly produced by cold isostatic pressing molding and densified by sintering at high temperature, but the sintered molybdenum is seriously brittle and has a large number of pores in the structure, which leads to low density and strength, and makes it difficult to meet the requirements of the application of high-temperature components. Highly deformed molybdenum is ductile both at room temperature and high temperatures, so the plastic deformation of molybdenum is particularly important [10]. The appropriate form of plastic deformation can improve the density of sintered molybdenum, reduce the DBTT [8, 11], improve the ductility, and greatly enhance its structural toughness and performance. In this article, we used the Gleeble 3500 thermo-mechanical simulator to study the influence of high-temperature compression on the microstructure and properties of sintered molybdenum through uniaxial compression experiment.

2 Experiment 2.1 Materials Cylindrical specimens of φ8 mm in diameter and 12 mm in height were taken from a commercial sintered molybdenum slab. The chemical composition is listed in Table 1, in which the Mo content is not lower than 99.95%. The density was 9.8 g cm−3 , which is determined by the Archimedes method. The average diameter of grains was about 31.8 μm, i.e., the number of grains per square millimeter was approximately 992.

Table 1 Chemical composition of sintered Mo

Element

ω/%

Element

ω/%

Fe

0.0017

P

≤0.0010

Ni

0.0005

C

0.0010

Al

≤0.0020

O

0.0030

Si

≤0.0030

N

≤0.0030

Ca

0.0005

W

0.0094

Mg

0.0003

Mo

≥99.95

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2.2 Experimental Method The Gleeble 3500 thermo-mechanical simulator was adopted, and a K-type thermocouple with the high-temperature cement protection was used to monitor and control the test temperature. Under the protection of high-purity argon atmosphere, the sintered Mo specimen was kept at 1200 °C for 3 min, and then the uniaxial compression was carried out under the condition of the deformation rate of 0.1 s−1 , with a deformation account of 50%. Before the experiment, nickel-based lubricant was coated on both ends of the anvil and specimen to reduce friction and improve deformation uniformity.

2.3 Measurements and Characterization The density of sintered and compressed molybdenum specimens was measured by using Archimedes method. The compression contact surface (A) and the longitudinal section (B) of the Mo specimen before and after the compression were analyzed by Vickers hardness, OM and XRD texture analysis, respectively. The specimen was first grounded, then mechanically polished until the surface was bright, and finally electropolished in 2 wt% NaOH aqueous solution at a voltage of 5–15 V for 10–20 s. The Vickers hardness was measured using a VTD 552 Vickers hardness tester. The load was 10 kg and the loading time was 10 s, and the description in this text was all represented by HV10/10. The microstructure of the Mo specimens was observed using an Axiovert Model 200MAT/Zeiss optical microscope. The corrosive agent was 10 wt% NaOH + 10 wt% Fe(KCN)3 aqueous solution. The X-ray diffraction texture data were collected on X’Pert MRD/Philips multifunction high-resolution diffractometer using a copper target with a voltage of 35 kV and a current of 20 mA. Table 2 is some of the symbols used in this paper and their implications.

Table 2 The symbol and implication

Symbol

Implication

Symbol

Implication

HV

Hardness (HV10/10)

S

The sintered Mo specimen

HVave

Average hardness

C

The compressed Mo specimen

HVper

Percentage of hardness fluctuation

A

Compressed contact surface

HVmax

Maximum hardness

B

Longitudinal section

HVmin

Minimum hardness

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3 Results and Discussion 3.1 Density As listed in Table 3, the size and density of the specimen before and after compression were measured, respectively. We found that after the high-temperature compression of the deformation amount of 50%, the barreling presented. The diameter of the end face and the barreling, respectively, increases by 30.6 and 49.7%, which indicates that the deformation is obviously inhomogeneous. The density of the specimen after compression is significantly increased, and reaches to ~99.4% of the theoretical density. This is mainly attributed to part of the sintered pores are flattened or even compacted during compression. And the densification could be raised by increasing the degree of deformation.

3.2 Vickers Hardness Table 4 shows the Vickers hardness results on each side of the sintered and compressed Mo specimens, and their average Vickers hardness (HVave ) and percentage of hardness fluctuation (HVper ) can be calculated from these hardness data. HVper is used to characterize the inhomogeneity of hardness distribution. It is simply defined by the following formula: HVper 

HVmax − HVmin × 100 HVave

(1)

From Table 4, it can be seen that the compressed contact surface (S-A) and the longitudinal section (S-B) of the sintered Mo specimen are, respectively, very close in terms of HVave and HVper , indicating that the hardness distribution of the specimen is homogeneous. But after compression at 1200 °C and 50% height reduction, the average Vickers hardness of the compression contact surface, and the longitudinal section of the compressed Mo specimen are significantly increased (C-A increased by 26.1%, C-B increased by 29.4%). Meanwhile, the percentage of hardness

Table 3 Size and density of sintered and compressed Mo specimens

Mo

Sintered Mo

Diameter/mm Height/mm Quality/g Density/g

7.9

10.32

11.98

6.02

6.215 cm−3

Relative density/%

Compressed Mo

6.118

9.8

10.16

95.9

99.4

11.83

Influence of High-Temperature Compression … Table 4 Vickers hardness of the sintered and compressed Mo specimens

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Sintered Mo

Compressed Mo

S-Aa

S-B

C-A

C-B

HVave

174.5

173.7

220.1

224.8

HVmax

183.1

180.5

235.8

237.7

HVmin

169.5

167.2

197.5

195.7

13.6

13.3

38.3

42.0

7.8

7.7

17.4

18.7

HVmax − HVmin HVper /% a In

this paper, S-A stands for the compression contact surface of sintered molybdenum, and the naming principles of S-B, C-A and C-B are similar

fluctuation gets larger (Sintered Mo: HVper  7.7–7.8%, Compressed Mo: HVper  17.4–18.7%), revealing that the compression deformation of Mo may be uneven. Comparing the compression contact surface and the longitudinal section of the compressed Mo specimen separately, it can also be found that the HVave and HVper of the longitudinal section are slightly higher than those of the compressed contact surface. Figure 1a–c is the hardness distribution curves of each side of the sintered Mo and compressed Mo, and Fig. 1d is a schematic diagram for measuring the hardness in the longitudinal section of compressed Mo. It can be seen from Fig. 1a that the hardness and hardness fluctuation of the sintered Mo specimen is relatively low. But the hardness of the Mo specimen is greatly improved after compression, as well as the hardness fluctuation, and the distribution shows a certain regularity. From Fig. 1b, it can be seen that on the contact surface of the sintered molybdenum specimen, the hardness in the center (S-A-center) and edge (S-A-edge) area are very close, and the hardness distribution are uniform. But for the compressed Mo specimen, the hardness of the compression contact surface is obviously improved; moreover, the hardness distribution shows the rule of low center and high edge with an evident difference. Figure 1c is the hardness distribution curve of the longitudinal section of compressed Mo specimen. As the curves illustrate, the hardness in this plane fluctuates greatly and the hardness distribution roughly shows the following rules: Hardness in the center area (C-B-center) is higher and more evenly distributed. Hardness in the left and right sides (C-B-left and right) is the second. And near the compression contact surface, hardness is the lowest in the up and down area (C-B-up and down).

3.3 Microstructure Figure 2 is the optical microstructure of the compression contact surfaces of sintered and compressed Mo specimens. Figure 2a shows a typical sintered structure. It can be seen from Fig. 2a that the grains of the sintered Mo are polygonal and equiaxed, the

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Fig. 1 a–c The hardness distribution curves of each side of the sintered Mo and compressed Mo, and d a schematic diagram for measuring the hardness of the compressed Mo longitudinal section

grain boundaries are clear and relatively flat. The grain size is not uniform and some grains are irregularly grown (indicated by arrows in Fig. 2a). Besides, there are a large number of dispersed pores remaining due to sintering shrinkage. The presence of these pores is the main reason why the density (9.8 g cm−3 ) of the sintered Mo specimen is lower than the theoretical density. It can be seen from Fig. 2b, c that on the compression contact surface of compressed Mo specimen, the microstructure between the center and edge area are obviously different. The center still maintains the equiaxed crystal structure similar to the sintered state. The grain boundary is relatively regular and clear. Compared with Fig. 2a, it can also be found that the number of pores at the grain boundary in Fig. 2b decreases, and the pores are mainly within the grains. In addition, the grain size increased obviously after compression. The number of crystals is about 496/mm2 , which is much lower than that 992/mm2 of the sintered Mo specimen. In Fig. 2c, the grains in the edge area are fractured, and a mass of substructures exist in more than half of the grains, and the grain boundaries are blurred (as shown by the arrow in Fig. 2c), and the pores almost disappear. As shown in Fig. 2d, there is also a sub-edge region between the center and the edge area (parallel to the interface line

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Fig. 2 Optical microstructure of sintered Mo specimen (a), and compressed Mo specimen at center (b), edge (c) and sub-edge (d)

shown in Fig. 2d). This is because, in the process of high-temperature compression, the central area of the compression contact surface is difficult to deform due to large frictional force. With the height reduction of the specimen, i.e., the deformation amount increased, the metal molybdenum located on the circumferential side of the cylindrical specimen is turned to the compression contact surface. The macro performance is the enlargement of the diameter (Table 3) and the barreling appears on the side of the circumferential. Some of the metal molybdenum grains that turn flat from the circumferential side to the compression contact surface are fractured, rotated and deformed, and the deformation degree of grain in this region is much larger than the grain in the central region of the compression contact surface. It can be concluded that the microstructure difference between the center and the (sub-edge) edge area of the compression contact surface is the reason why the hardness presents the rule of low center and high edge. Figure 3 shows the metallographic microstructure of the longitudinal section of the compressed Mo specimen and the observation position of each metallographic micrograph. Figure 3a, b represent the optical microstructure of the upper and lower regions of the longitudinal section. We found that near the compression contact surface, the deformation degree of grain is small, the grain boundaries are clear, and there are still plenty of pores in the structure. Compared with Fig. 2b, the large-diameter grains with

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Fig. 3 Optical microstructure of the longitudinal section of the compressed Mo specimen (a–e) and observation position (f). a Upper part; b lower part; c center; d upper left; e upper right

slight elongation are generated, while some grains are fractured and the grain size uniformity is reduced. As shown in Fig. 3d, in the longitudinal section, comparing the marginal and middle part near the compression contact surface, it is found that the grain deformation of the marginal part is more sufficient than that in the middle region and generates a distinct intermediate zone (as the line shown in Fig. 3d). In combination with Fig. 2d, it can be found that close to the center of the compression contact surface, a region which is nearly “hemispherical” and whose deformation degree of grains is significantly lower than other areas is formed. The reason for the formation of this region is that during the compression, there is strong friction

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between the anvil and the specimen, which prevents the grains in this region to slide and deform. And this is also the reason why the hardness of this region is lower than other locations. This region is often referred to as the hard-deforming zone (zone I) [8], and in some other reports, it is also called dead metal zone (DMZ) [12, 13]. Figure 3c shows that the grain deformation is most pronounced in the central region of the longitudinal section. Perpendicular to the compression direction, the grains are flattened and elongated to generate a curved “laminate” deformed structure containing abundant substructures. This is attributed to the grains in this area are least affected by friction and subjected to three-direction compressive stress. The flow of metal particles is more likely to happen and the microstructure deforms most sufficiently. Therefore, the hardness is the highest in the central region which belongs to the easy-deforming zone (zone II) [8]. Figure 3e shows optical microstructure at the edge of the upper right corner, corresponding to the part of the specimen in Fig. 2c where the circumferential metal is flattened to the end faces during compression. It is considered that the grains in this region are subjected to the extrusion of zone I along the radius, and the deformation occurs when the grains are flattened from the circumferential side to the end face. At the same time, the grains are influenced by a certain amount of friction from the anvil and other grains. Therefore, the grain deformation degree of this area is between zone II and zone I, and the hardness is also between the two. This region is alluded to as the free-deforming zone (zone III) [8]. In summary, because of the difference in the deformation degree of various regions, the changes of the density of grain boundaries and the substructure, as well as the number and size of pores, result in a certain regular distribution of the Vickers hardness of the compressed Mo specimen. Although the overall properties of the compressed Mo such as hardness, density, and strength are greatly improved compared with the sintered Mo, there is still much room for improvement in the homogeneity of high-temperature compression.

3.4 Textures Figure 4 is the orientation distribution function (ϕ2  45 ◦ ) before and after the compression of sintered Mo. It is found in Fig. 4a that there is a 113//ND fiber texture with a maximum intensity of 9.478 in the sintered Mo. This sintering texture should be the result of minimizing the energy of grain boundaries. During the sintering process, in order to minimize the boundary energy between grains, the randomly distributed grains will rotate toward the lowest energy position, leading to the aggregation of the orientation corresponding to the low-energy boundary, and finally generating the sintered texture [14]. As shown in Fig. 4b, c, the texture components of Mo specimens changed significantly after compression, and different types of textures were observed in compression contact surfaces and longitudinal sections. The compression contact surface

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Fig. 4 The ϕ2  45 ◦ ODF map of a sintered Mo, maximum intensity  9.478; b compression contact surface of the compressed Mo, maximum intensity  5.302; c longitudinal section of the compressed Mo, intensitymax  6.949; d main texture components of rolled BCC metal. Reprinted from Oertel et al. [15], Copyright 2010, with permission from Elsevier

mainly consists of the texture of γ-fiber, i.e., 111//ND, with the maximum intensity of 5.302. In the longitudinal section, four types of textures were found, which are {112}111 copper texture, {001}100 cubic texture, {001}110 rotation cube texture, and residual {113}111 sintered texture with the highest intensity of 6.949. There are two possible reasons for the above complex deformed texture: First, the severe uneven deformation leads to the formation of different textures in different regions. Second, due to insufficient deformation, many textures have not yet reached the optimal orientation, such as residual {113}111 sintered texture. We also found that the texture intensity of the sintered Mo is stronger than that of the compressed Mo. During the high-temperature compression process, the

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deformation causes deformed textures in the metal Mo specimen, and the dynamic recovery and recrystallization lead to the recrystallized textures in the specimen, and the two processes are simultaneously or alternately performed. Both types of textures could not be fully developed, which leads to texture weakening after high-temperature compression. In addition, the deformation of sintered Mo could be affected by many factors such as deformation amount, deformation temperature, strain rate, initial grain size, and so on. The deformed textures and recrystallized textures will be promoted or suppressed in varying degrees, which will influence the type and intensity of textures after deformation [16].

4 Conclusion (1) After the high-temperature compression of the sintered Mo at 1200 °C, strain rate of 0.1 s−1 , and deformation amount of 50%, a large number of internal pores are flattened and compacted, the relative density is increased to nearly theoretical density. (2) The average Vickers hardness increases from 174.5 to 224.1 MPa after hightemperature compression. However, the homogeneity of the compressed Mo is reduced compared with that of sintered Mo. (3) The uniaxial compression of sintered Mo generates a hard-deforming zone (zone I), a free-deforming zone (zone III), and an easy-deforming zone (zone II). The deformation degree of zone I is the lowest, and it remains nearly the assintered structure. The grains in zone II are deformed most sufficiently, forming a “lamellar” structure containing plenty of substructures, showing the largest deformation degree. (4) There is a certain intensity of 113//ND fiber texture in the sintered Mo. 111//ND fiber texture, {112}111 copper texture, {001}100 cube texture, {001}110 rotating cube texture, and {113}111 residual sintered texture are observed in the compressed Mo. Acknowledgements This work was supported by the National Key Research and Development Program of China (Grant No. 2017YFB0306000).

References M. Scapin, L. Peroni, F. Carra, Investigation and mechanical modelling of pure molybdenum at high strain-rate and temperature. J. Dyn. Behav. Mater. 2(4), 460–475 (2016) G. Liu, G.J. Zhang, F. Jiang et al., Nanostructured high-strength molybdenum alloys with unprecedented tensile ductility. Nat. Mater. 12(4), 344–350 (2013) C.G. Oertel, I. Huensche, W. Skrotzki et al., Plastic anisotropy of straight and cross rolled molybdenum sheets. Mater. Sci. Eng., A 483(1), 79–83 (2008)

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S. Primig, H. Leitner, H. Clemens et al., On the recrystallization behavior of technically pure molybdenum. Int. J. Refract Metal Hard Mater. 28(6), 703–708 (2010) K. Babinsky, S. Primig, W. Knabl et al., Fracture behavior and delamination toughening of molybdenum in Charpy impact tests. JOM 68(11), 2854–2863 (2016) J. Wadsworth, C.M. Packer, P.M. Chewey et al., A microstructural investigation of the origin of brittle behavior in the transverse direction in Mo-based alloy bars. Metall. Trans. A 15(9), 1741–1752 (1984) A. Kumar, B.L. Eyre, Grain boundary segregation and intergranular fracture in molybdenum. Proc. R. Soc. London 370(1743), 431–458 (1980) P. Zhihui, Rare Metal Material Processing Technology. Central South University Press, Changsha (2003) C. Chen, H.Q. Yin, X.H. Qu et al., Research of deformation resistance of molybdenum. Rare Metal Mater. Eng. 36(7), 1237–1240 (2007) J.W. Christian, Some surprising features of the plastic deformation of body-centered cubic metals and alloys. Metall. Trans. A 14(7), 1237–1256 (1983) V.P. Pilyugin, L.M. Voronova, T.M. Gapontseva et al., Structure and hardness of molybdenum upon deformation under pressure at room and cryogenic temperatures. Int. J. Refract Metal Hard Mater. 43(12), 59–63 (2014) L. Zaharia, R. Comaneci, R. Chelariu et al., A new severe plastic deformation method by repetitive extrusion and upsetting. Mater. Sci. Eng., A 595(5), 135–142 (2014) E. Tempelman, B.N.V. Eyben, H. Shercliff, Manufacturing and design (Butterworth-Heinemann, Oxford, 2014) G. Herrmann, H. Gleiter, G. Bäro, Investigation of low energy grain boundaries in metals by a sintering technique. Acta Metall. 24(4), 353–359 (1976) C.G. Oertel, I. Hünsche, W. Skrotzki et al., Influence of cross rolling and heat treatment on texture and forming properties of molybdenum sheets. Int. J. Refract Metal Hard Mater. 28(6), 722–727 (2010) M. Weimin, Crystallographic Texture and Anisotropy of Metallic Materials. Science Press, Beijing (2002)

Microstructure and Mechanical Properties of a Ti–Al–Sn–Zr–Mo–Nb–W–Si High-Temperature Titanium Alloy Y. W. Diao, X. Y. Song, W. J. Zhang, M. Y. Zhao, W. J. Ye and S. X. Hui

Abstract The Ti–6.5Al–2Sn–4Zr–1Mo–3Nb–0.5W–0.2Si (1M3N0.5W) alloy is a novel two-phase high-temperature alloy for short-term use up to 700 °C. The effects of different heat treatment regimes on the microstructure and mechanical properties were investigated through optical microscopy (OM), scanning electron microscopy (SEM), and tensile tests at the temperatures up to 700 °C. The results show that 1M3N0.5W alloy treated after single-stage annealing can get bimodal structure, which consists of equiaxed primary α (αp ) phase and lamellar transformed β (βt ) structure. The ultimate tensile strength (UTS) and 0.2% yield strength (YS) of the alloy at 700 °C are 473 and 310 MPa, respectively. The 1M3N0.5W alloy treated after solution and aging treatment can get Widmanstatten structure, which consists of coarse primary β-grain and secondary α-phase precipitated on the β-grain. Compared with single-stage annealed alloy, the UTS and YS of the solution and aging treated alloy are significantly improved at 700 °C, while the room-temperature plasticity is dramatically reduced.

1 Introduction High-temperature titanium alloys are widely used in aerospace industry such as compressor and air-wing of aero-engine owing to their high strength to weight ratio, hightemperature creep resistance, fatigue strength, endurance strength, and structure stability [1–6]. So far, the current temperature limit of the conventional high-temperature titanium alloys is about 600 °C. Several high-performance titanium alloys have been developed such as IMI834 [7], Ti1100 [8], BT36 [9], and Ti-60 [10]. With the development of the aerospace industry, sky-rocketing development of high-speed vehicle has intensified the research on improving the material property. Traditional long-term high-temperature titanium alloys is not able to satisfy the need of use [11, 12]. Generally, different from long-term high-temperature Y. W. Diao · X. Y. Song (B) · W. J. Zhang · M. Y. Zhao · W. J. Ye · S. X. Hui State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co. Ltd., Beijing 101407, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_9

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titanium alloy, the service temperature of short-term high-temperature titanium alloy can reach 650 °C even 700 °C, and the endurance performance under large stress (close to yield strength) is better. With the increase in temperature, the mechanical properties of high-temperature alloys will be significantly changed. In consideration of mechanical properties, forming process, etc., the research work of short-term high-temperature titanium alloy used above 650 °C has received great attention. In titanium alloy, α + β two-phase alloy shows good high-temperature performance and processing performance [13]. Therefore, on the basis of the traditional Ti–Al–Sn–Zr–Mo–Si alloy system, our research group has added Nb and W elements to improve the high-temperature strength and workability of the alloy, and designed the new Ti–6.5Al–2Sn–4Zr–Mo–Nb–W–0.2Si series α + β two-phase titanium alloys. The alloys show excellent processing performance and can be shortly applied under the environment of 700 °C. In this paper, Ti–6.5Al–2Sn–4Zr–1Mo–3Nb–0.5W–0.2Si alloy was taken as the research object. Two different heat treatments were adopted, and the effect of heat treatment on the microstructure and mechanical properties of the alloy was systematically studied.

2 Experimental The test material was Ti–6.5Al–2Sn–4Zr–1Mo–3Nb–0.5W–0.2Si (hereinafter termed as 1M3N0.5W) alloy. The ingot was melted three times using vacuum consumable electrode method, and then forged and hot-rolled into 12 mm-diameter bars in the two-phase region. The β-transus temperature (T β ) of this alloy was measured to be 1110–1020 °C. Two different kinds of heat treatments were, respectively, carried out for the alloy, i.e. 990 °C/2 h/AC (single-stage annealing) and 1040 °C/1 h/AC + 600 °C/2 h/AC (solution and aging treatment). For OM and SEM observation, the specimens were grounded to 5000# SiC paper and then etched by a solution of 30 ml H2 O, 2 ml HF and 6 ml HNO3 . The microstructures under different conditions were observed using an Axiovert 200 MAT type Zeiss optical microscope and JSM-6510 type scanning electron microscope. Tensile tests were performed in air at room temperature and 700 °C using an MTS universal tensile testing machine. The tensile specimens had gauge dimensions of 5 mm in diameter and 25 mm in length. Three tensile specimens were tested to get an average.

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3 Results and Discussion 3.1 Microstructures After Different Heat Treatments Generally, the alloy deforms in the upper part of the two-phase region to obtain a bimodal structure. Figure 1 shows the microstructure of 1M3N0.5W alloy after single-stage annealing treatment. It can be seen that after single-stage annealing, a bimodal structure composed of equiaxed primary α (αp ) phase and lamellar transformed β (βt ) structure is obtained. The αp phase content is about 40%, and the secondary α phase (αs ) with a size of about 300 nm is distributed on the lamellar βt structure. The alloy was heat treated in the β single-phase region to obtain a Widmanstatten structure. Figure 2 shows the microstructure of 1M3N0.5W alloy after (T β + 20 °C)/1 h /AC + 600 °C/2 h/AC heat treatment. It can be seen that the alloy is composed of Widmanstatten structure after solution and aging treatment, which consists of coarse raw β-grain and αs phase precipitated on the β-grain. Due to the highsolution temperature above T β , the grains get full growth, so the β grain boundaries are clear and intact. The size of α lath bundles of 1M3N0.5W alloy after treatment at (T β + 20 °C)/1 h/AC + 600 °C/2 °C/AC is approximately 400–1000 nm.

3.2 Mechanical Properties After Different Heat Treatment Figure 3 shows the tensile properties of 1M3N0.5W alloy heat treated at different temperatures after (T β -20 °C)/2 h/AC single-stage annealing, including UTS, YS, and elongation (EL). It can be seen that the values of UTS, YS, and EL of the alloy at room temperature are, respectively, 1140 MPa, 1050 MPa, and 16%. With the increase of tensile temperature, the UTS and YS of the alloy are decreased significantly and the EL is obviously increased. The average UTS and YS of the tested alloy at 700 °C are

Fig. 1 Microstructure of 1M3N0.5W alloy heat-treated at 990 °C/2 h, AC: a OM and b SEM

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Fig. 2 Microstructures of 1M3N0.5W alloy heat-treated at 1040 °C/1 h/AC + 600 °C/2 h/AC: a OM and b SEM Fig. 3 Tensile properties as a function of testing temperature for alloy after single-stage annealing

473 and 310 MPa, respectively, and the EL can reach 52%. Comparing the changes of UTS and YS at temperatures from 600 to 700 °C, it can be noted that the change of YS at high temperature is more sensitive than that of UTS. The tensile properties of 1M3N0.5W alloy after solution and aging treatment are tested at room temperature and 700 °C are listed in Table 1. Similar to the alloy after single-stage annealing, the UTS and YS are decreased as temperatures increases while EL is increased. The room-temperature tensile properties of 1M3N0.5 W alloy after solution and aging treatment are compared with that after single-stage annealing treatment, as shown in Fig. 4. It is found that compared with single-stage annealing, the UTS, YS, and EL of the alloy at room temperature are reduced after solution and aging treatment. Studies have shown that the room-temperature plasticity of 1M3N0.5W alloy with Widmanstatten structure is significantly lower than that of the bimodal structure. According to the previous research, the reason can be contributed to that the microstructure of the bimodal structure is smaller than that of the Widmanstatten structure, which will make it easier to generate coordination deformation. In the deformation behavior of titanium alloy, when the alloy is solution

Microstructure and Mechanical Properties … Table 1 Tensile properties of the alloy after solution and aging treatment at different temperatures

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Temperature, °C

YS, MPa

UTS, MPa

EL, %

Room temperature

956

1119

6

700

421

581

25

Fig. 4 Comparison of tensile properties at room temperature under different heat treatments

treated in the β-phase region, the β-phase will be easy to growth which will cause coarse grains, and coarse grains can reduce the room-temperature plasticity of the alloy. Generally, the Widmanstatten structure may have good creep properties, hightemperature strength, and fracture toughness for high-temperature titanium alloys [14], its oxidation resistance is also significantly better than that of bimodal structure [15]. The presence of the αs phase reduces the intergranular fracture ratio, different orientations of the α lath bundles make the crack expand along the α/β phase interface, their combined effect is beneficial to improve the toughness and tensile properties of the alloy. However, due to the inevitable β-grain growth and the formation of the continuous network-like α-grain boundary, the coordination deformation ability reduced, and its elongation reduced accordingly. Figure 5 shows the comparison of the tensile properties of single-stage annealed alloy and solution and aging treated alloy at 700 °C. According to Fig. 5, after solution and aging treatment, the values of UTS and YS are 581 MPa and 421 MPa at 700 °C, respectively, and value of EL can reach 25%. It can be seen from Fig. 5 that compared with single-stage annealed alloy, the UTS and YS of the solution and aging treated alloy have been significantly improved at 700 °C. The UTS and YS are increased by 22 and 30%, respectively. Studies have shown that the high-temperature strength of 1M3N0.5W alloy of Widmanstatten microstructure is significantly higher than that of bimodal structure, but its EL is lower. Its high-temperature strength is mainly due to the αs phase which grows in the β-grain constitutes a lamellar structure, and there are many short needle-like αs

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Fig. 5 Comparison of tensile properties at 700 °C under different heat treatments

phases formed among these parallel lamellar structures. The two types of αs phase work together and result in high strengthening effect. Microstructure has a strong influence on the mechanical properties of titanium alloy. The choice of structure is usually determined by the performance requirements of the alloy in the environment of use. As for the short-term application of hightemperature titanium alloy, it is hoped that the alloy will get good room-temperature ductility and workability while ensuring the high-temperature strength of the alloy. From the paper, we can know that the bimodal structure exhibits a better combination of high-temperature mechanical properties and room-temperature ductility. So it can be proved that for the alloy, the single-stage annealing (990 °C/2 h/AC) has greater potential for application.

4 Conclusions In this present work, the microstructure and mechanical properties of Ti–6.5Al–2Sn–4Zr–1Mo–3Nb–0.5W–0.2Si alloy with different heat treatments have been investigated in detail, and the following conclusions can be drawn: (1) After single-stage annealing, 1M3N0.5W alloy can get bimodal structure, consisting of equiaxed αp phase and lamellar βt structure. After solution and aging treatment, the alloy can get Widmannstatten structure which consists of coarse raw β-grain and secondary α-phase precipitated on the β-grain. (2) The average UTS and YS at 700 °C of the alloy after single-stage annealing are 473 and 310 MPa, respectively. With increasing testing temperature, the change of YS is more sensitive than that of UTS.

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(3) The room-temperature plasticity of Widmannstatten structure is significantly lower than that of the bimodal structure. The high-temperature strength is significantly higher than that of bimodal structure, but its EL is lower.

References 1. T. Li, et al., Effect of heat treatment processes on microstructure and properties of TC4-DT titanium alloy. Titan. Ind. Prog. (2017) 2. J. Cheng, C. Shi., Study progress of microstructure, properties and processing technology of titanium alloys. Hot Work. Technol. (2016) 3. M. Jackson, K. Dring, A review of advances in processing and metallurgy of titanium alloys. Metal Sci. J. 22(8), 881–887 (2006) 4. Z. Lian., Review of titanium industry progress in America, Japan and China. Rare Metal Mater. Eng. 32(8), x-584 (2003) 5. Y. Zhao et al., The high temperature deformation behavior and microstructure of TC21 titanium alloy. Mater. Sci. Eng., A 527(21), 5360–5367 (2010) 6. J.F. Lei et al., LFatigue crack initiation in Ti–5Al–4Sn–2Zr–1Mo–0.7Nd–0.25Si high temperature titanium alloy. Int. J. Fatigue 19(93), 95–98 (1997) 7. J. Mallol, M.C. Sarraga, M. Bartolomé et al., Hot working behavior of near-α alloy IMI834. Mater. Sci. Eng. A 396(1), 50–60 (2005) 8. F. Ma et al., Mechanical properties and strengthening effects of in situ, (TiB+TiC)/Ti-1100 composite at elevated temperatures. Mater. Sci. Eng. A 654, 352–358 (2016) 9. Mengyi Hao, J. Cai, D.U. Juan, The effect of heat treatment on microstructure and properties of BT36 high temperature alloy. J. Aeronaut. Mater. 23(2), 14–17 (2003) 10. Weiju Jia et al., High-temperature deformation behavior of Ti60 titanium alloy. Mater. Sci. Eng. A 528(12), 4068–4074 (2011) 11. Y. Wang, B. Lu, R. Yang, Effect of (Mo, W) content and heat treatment on the tensile properties of a high-temperature titanium alloy. Chin. J. Mater. Res. 24(3), 283–288 (2010) 12. Y. Zhang, D. Guo, et al., Growth behavior of α phase in Ti-5.6Al-4.8Sn-2.0Zr-1.0Mo-0.35Si0.7Nd titanium alloy. Mater. Sci. Technol. 22(4), 459–464 (2006) 13. X. Gong, et al., Experimental studies on the dynamic tensile behavior of Ti–6Al–2Sn–2Zr–3Mo–1Cr–2Nb–Si alloy with Widmanstatten microstructure at elevated temperatures. Mater. Sci. Eng. A 523(1), 53–59 (2009) 14. C Leyens, M Peters, Titanium and titanium alloys: fundamentals and applications. Titanium and Titanium Alloys (2003) 15. K.V.S. Srinadh, V. Singh, Oxidation behaviour of the near α-titanium alloy IMI 834. Bull. Mater. Sci. 27(4), 347–354 (2004)

Optimization of Cold-Rolling-Stabilization Process for High Mg-Containing Al Alloy Yueying Liang, Hui Huang, Xiaolan Wu, Shengping Wen, Kunyuan Gao, Zuoren Nie and Hongmei Li

Abstract The mechanical properties and corrosion properties of 5E61 alloy coldrolled sheet were studied under different stabilization annealing processes. The microstructure after alloy stabilization annealing was observed by optical microscopy (OM), scanning electron microscopy (SEM) and transmission electron microscopy (TEM), and the cold-rolled sheet stabilization process window was obtained. The results show that the 5E61 alloy has good corrosion resistance under different stabilization annealing processes; High strength and good corrosion resistance can be achieved in a short time in a higher stabilizing annealing temperature range; while after treatment in the lower stabilized annealing temperature range, the results show that the corrosion resistance is good under long-term annealing conditions and high mechanical properties are obtained, thereby the optimization of the stabilization process was carried out.

Y. Liang (B) · H. Huang · X. Wu · S. Wen · K. Gao · Z. Nie · H. Li School of Materials Science and Engineering, Beijing University of Technology, Beijing 100124, People’s Republic of China e-mail: [email protected] H. Huang e-mail: [email protected] X. Wu e-mail: [email protected] S. Wen e-mail: [email protected] K. Gao e-mail: [email protected] Z. Nie e-mail: [email protected] H. Li e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_10

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1 Introduction Al–Mg aluminum alloy has high specific strength, low density, good weldability, and high strength and corrosion resistance in the marine environment. Therefore, 5xxx series aluminum alloy is widely used in structural materials of marine and civil ship structures [1, 2]. Since the aluminum alloy of this series is not heat-treated, the strength of the alloy is mainly improved by solid solution strengthening and cold deformation, which limits its performance and application. Studies have shown that [3]: The microalloying method can improve the comprehensive performance of the 5xxx series aluminum alloy. The addition of trace element Er can form a fine Al3 Er dispersion phase in the alloy matrix, and has good interface matching with α-Al, which produces strong pinning dislocations, resulting in precipitation strengthening effect [4]. It is found that the sensitization phenomenon is related to the continuous formation of the β phase (Al3 Mg2 ) along the grain boundary, which reduces the corrosion resistance of the alloy [5]; and the content of β phase in high Mg-containing aluminum alloy varies greatly with temperature [6]. In this paper, the mechanical properties and corrosion properties of Er-high magnesium alloy 5E61 cold-rolled sheet under different stabilization annealing processes were studied, and the microstructural analysis was carried out on typical conditions to further optimize the stabilization process of the alloy.

2 Experimental 2.1 Material Preparation The initial material used in the test was a 7 mm thick plate of 5E61-H112 state, and its chemical composition was measured by X-ray fluorescence spectroscopy (XRF) as listed in Table 1. The 5E61-H112 aluminum alloy was placed in an air circulating heating furnace for intermediate annealing at 350 °C/2 h, then air cooled to room temperature, then cold rolled in multiple passes, each pass was controlled to 10–25% reduction, the total cold reduction was 50%, and finally the 5E61-H16 alloy was obtained. The cold-rolled alloys were annealed at 235, 250 °C/(1, 2, 4, 8, 16, 24 h), 270 °C/(1–10 h), 285 °C/(1–10 h), respectively. The alloys in the corresponding state were subjected to room temperature tensile test, intergranular corrosion test, exfoliation corrosion test, metallographic structure, and transmission electron microscopy (TEM) observation and analysis.

Table 1 Experimental alloy composition (%, mass fraction)

Mg Mn Er 5E61

Zr

Fe

Si

Cu Zn Al

5.75 0.84 0.20 0.09 0.13 0.09 0.05 0.02 Bal.

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2.2 Characterization The hardness of the alloy under different stabilization annealing treatments of 5E61 was tested by HXD-1000 microhardness tester. The load size was 1000 gf, the loading time was 10 s, and each sample was randomly selected after 10 points. Samples for tensile tests were prepared according to the processing standard of GB-T16865. The experiment was carried out on an MTS810 material testing machine with a tensile rate of 1 mm/min. Three samples for each state were tested, and the results were averaged. The intergranular corrosion test was carried out in accordance with ASTM G67, and the intergranular corrosion sensitivity of the material was evaluated by the mass loss per unit area after the intergranular corrosion test. The sample size of the experiment is 50 mm * 6 mm * 3.5 mm. Before the experiment, the surface of the sample was polished with 320# sandpaper; washed, immersed in 65–68 wt% concentrated HNO3 solution for 30 s, washed with water, dried; the sample was tilted to the bottom and side walls of the container (capped), and concentrated HNO3 (5 ml/cm2 ) was added, etching in a constant temperature water bath environment of 30 °C for 24 h; the corroded sample was taken out and washed with deionized water and air-dried; weighed, and finally the weight loss per unit area was expressed as mg/cm2 . The peeling corrosion test was carried out in accordance with ASTM G66-99 (2013). The sample was washed with acetone, etched in a 5 wt% NaOH solution at 80 °C for 1 min, washed with water, decontaminated with 65–68 wt% concentrated HNO3 for 30 s, washed with deionized water, and dried; The etching solution was prepared from NH4 Cl 53.5 g, NH4 NO3 20.0 g, (NH4 )2 C4 H4 O6 1.8 g, 30% H2 O2 10 ml and deionized water to prepare a 1 L solution. The PH of the solution was finally adjusted to 5.36. The sample was immersed at least 25 mm below the liquid surface with a cotton thread, 25 mm from the bottom of the container, and the container was capped and kept at a constant temperature of (65 ± 1)°C for 24 h. The sample after corrosion is taken out and rinsed rapidly, and the concentrated HNO3 is immersed in the room temperature to remove the surface, then washed, and dried. Finally, the sensitivity to exfoliation corrosion was evaluated by macroscopic testing standards. Corrosion samples in the sensitized state after sensitization of different stabilization annealing processes were observed and analyzed by optical microscopy. The samples were polished by 400#, 800#, 1500#, 2000# sandpaper, and after polishing, the corrosion depth was observed by a BX51M optical microscope. Thin foils for TEM observation were prepared by the electrolytic double jet method. The electrolyte was a 30% concentrated HNO3 + 70% CH3 OH (vol.%) mixed solution. The sample was cut from a sheet of about 2 mm, and then mechanically grounded to 80–120 μm, and then punched into a wafer having a diameter of about 3 mm, electrolytic double spray thinning on the Struers Tenupol-2 dual sprayer, voltage around 12–15 V, current 80–100 mA, the temperature is controlled at around −30 °C. Sample observations were performed on a JEOL 2010 transmission electron microscope with an operating voltage of 200 kV.

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3 Results and Discussion 3.1 Microhardness Figure 1 is a microhardness of a 5E61-H16 alloy after annealing at 500 °C for 1 h. It can be seen that: as the annealing temperature increases, the microhardness decreases, where the hardness of the alloy decreases sharply between 250 and 300 °C. The 50% cold-deformed 5E61 alloy has recrystallization starting temperature of 250 °C and a final temperature of 300 °C; for this alloy, the sensitization temperature of the corrosion performance is generally within the recovery temperature range. Therefore, according to the hardness change curve, 235, 250 °C/(1, 2, 4, 8, 16, 24 h) and 270, 285 °C/(1, 2, 4, 6, 8, 10 h) were selected as the stabilization annealing process explored herein.

3.2 Tensile Properties Figure 2 is the tensile properties of the initial state and different annealed states after 50% cold-rolling deformation. It can be seen from the tensile strength and yield strength of Fig. 2a, b, the strength decreases rapidly after annealing for 1 h at different temperatures, especially at 270 and 285 °C, the severity of the alloy was significantly higher than 235 and 250 °C. After that, the alloy strength tends to decrease as the annealing time increases. It can be seen from Fig. 2c that the elongation of the experimental alloy rises rapidly within 1 h when annealed at different temperatures, and the increment after annealing at 270 and 285 °C is significantly higher than that at 235 and 250 °C. At the same annealing time, the elongation of the experimental alloy increase with the increase of annealing temperature. The strength of the alloy

Fig. 1 The change of microhardness after annealing for 1 h at different temperatures

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Fig. 2 Tensile properties of 5E61 alloy after different stabilization annealing processes

after annealing at two relatively low temperatures of 235 and 250 °C is significantly higher than the other two higher annealing temperatures, then the degree of decline in alloy strength tends to be stable over time. This is attributed to the Er and Zr elements are more favorable for the precipitation of the Al3 (Er, Zr) second phase particles in a relatively low temperature range, so that the inside of the alloy can pin the grain boundaries and dislocations, inhibition of recrystallization, thereby improving the mechanical properties of the alloy.

3.3 Long-Term Corrosion Performance of 5E61 Alloy in Different Stabilized Annealed States In order to investigate long-term intergranular corrosion resistance of 5E61 alloy in different annealed states, the intergranular corrosion test before and after sensitization at 100 °C/168 h was carried out, respectively. Figure 3 shows the alloy isothermally annealed at 235 and 250 °C. A graph of the weight loss value as a function of time before and after sensitization. It can be seen from the figure that the intergranular corrosion insensitive zone is entered after 4 h in the isothermal annealing

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Fig. 3 Curve of weight loss versus time of experimental alloys before and after sensitization at 235 and 250 °C

process at 235 °C, further extending the annealing time, the weight loss value remains basically unchanged; during the isothermal annealing at 250 °C, the alloy enters the intergranular corrosion insensitive zone after 2 h and remains essentially unchanged. Further analysis of the weight loss curves after sensitization of different annealed states shows that: After 235 °C isothermal annealing for 4 h, the sensitized alloy enters the medium sensitive region from the severely corroded zone, with the extension of annealing time, the weight loss after annealing for 8 h is 11.0 mg/cm2 , and enters the insensitive zone; the alloy with isothermal annealing and sensitization at 250 °C decreased rapidly with the extension of annealing time, and entered the intergranular corrosion insensitive region at 8 h, but the alloy entered the sensitive region with the annealing time. In order to further explore the intergranular corrosion of the alloy in the medium-sensitive region, the OM observation of the alloy in the sensitized state after sensitization is carried out, as shown in Fig. 4a–d, the alloy enters the passivation zone at 250 °C/4 h. It can be seen from the OM photograph of this state that the alloy has slight pitting corrosion and does not enter the alloy to cause intergranular corrosion. A slight pitting occurred when annealing at 235 °C for 8 h, the annealing time was further extended to 16 and 24 h, it can be seen that only slight pitting occurred and no intergranular corrosion occurred. At this time, the intergranular corrosion resistance of the alloy was good. The above analysis shows that the annealing resistance of the alloy at different temperatures is gradually improved with time. To investigate the effect of higher stabilization annealing temperature on the properties of 5E61 alloy, Fig. 5 is a graph showing the change of weight loss with time before and after sensitization of the experimental alloy at 270 and 285 °C. It can be seen from the figure that during the isothermal annealing process at 270 and 285 °C, the weight loss value of the alloy remains basically unchanged with the extension of the annealing time, and both are in the intergranular corrosion insensitive zone, indicating that the alloy has good resistance to intergranular corrosion under the stabilization annealing process. Similarly, in order to explore the long-term resistance

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Fig. 4 Optical micrographs of the 5E61 alloy in the medium sensitized state after isothermal annealing sensitization at 235 and 250 °C Fig. 5 Curve of weight loss with time before and after sensitization of the experimental alloy at 270 and 285 °C

of the experimental alloy to intergranular corrosion, the cold-rolled sheets treated with different annealed states were sensitized, the stabilizing annealing process of the alloy was further optimized by analyzing the intergranular corrosion properties after sensitization. It can be seen from the figure that the weight loss value of the two alloys after sensitization is gradually decreased with the extension of annealing time, and they all enter the medium sensitive area after annealing for 2 h; With the increase of annealing time, the alloys of the two annealing processes entered the insensitive zone at 6 h. Combined with Fig. 6, it can be seen that when pitting at 270 °C for 2 h, slight pitting occurred, no intergranular corrosion occurred, and the intergranular corrosion resistance of the alloy was good; when the alloy was annealed for 4 h, the alloy did not undergo pitting and intergranular corrosion. At this time, the intergranular corrosion resistance of the experimental alloy was good. Pitting and intergranular corrosion did not occur in 285 °C isothermal annealing for 2–4 h. At this time, the intergranular corrosion resistance was better. The above analysis shows that the long-term anti-intergranular corrosion resistance gradually becomes better as the annealing time is prolonged, and the results after treatment with the above two relatively low stabilization annealing temperatures show that the experimental alloy can obtain good resistance to intergranular corrosion after annealing in a short time.

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Fig. 6 Optical microscopic observation of the 5E61 alloy in the medium sensitized state after isothermal annealing sensitization at 235 and 250 °C

3.4 Exfoliation Corrosion Performance Test Exfoliation corrosion test was carried out on the 5E61 alloy cold-rolled sheet after different stabilization annealing treatment. Figures 7, 8, 9 and 10 show the macroscopic morphology of exfoliation corrosion in cold-rolled state and different annealed states, respectively. According to the standard evaluation, the cold-rolled alloy peeling corrosion grade in Fig. 7 is EB grade. After annealing at 235 °C/1 h, the alloy showed slight delamination and pitting and bubbling on the surface, rating EA. After annealing at 235 °C/2 h, the surface of the alloy was pitting, and the rating was PB. When the annealing time was 4–24 h, no pitting and delamination occurred on the surface, and the rating was N grade. Figure 8 shows that annealing at 250 °C for 1–4 h, the surface of the alloy only slightly pitting, rated as PA grade. After further annealing to 24 h, no pitting and delamination occurred on the surface of the alloy, and the rating was N. Figure 9 shows the macroscopic morphology after isothermal annealing at 270 °C, the alloy only showed slight pitting at 1–8 h, and the rating was PA grade; no pitting and delamination occurred on the alloy surface at 270 °C/10 h, and the rating was N grade. The macroscopic morphology of exfoliation corrosion after isothermal annealing at 285 °C shows that pitting and delamination have not occurred in the process of 1–10 h, and the grade is N grade, which has good anti-flaking corrosion performance. The above analysis shows that the experimental alloy is isothermally annealed at different times, and the anti-flaking corrosion performance gradually increases with time. The increase of the stabilization annealing temperature is beneficial to improve the anti-flaking corrosion performance of the alloy.

3.5 TEM Observation The corrosion performance of Al–Mg alloys is greatly affected by the distribution of β phase. Studies have shown that dislocation networks, grain boundaries and second phase particles preexisting inside the matrix contribute to β phase precipitation [7]. At the same time, Gaies et al. [8] found that the β phase is more inclined to precipitate

Optimization of Cold-Rolling-Stabilization Process … Fig. 7 Macroscopic morphology of exfoliation corrosion of 5E61 alloy after cold-rolling and isothermal annealing at 235 °C

Fig. 8 Macroscopic morphology of exfoliation corrosion of 5E61 alloy after isothermal annealing at 250 °C

Fig. 9 Macroscopic morphology of exfoliation corrosion of 5E61 alloy after isothermal annealing at 270 °C

Fig. 10 Macroscopic morphology of exfoliation corrosion of 5E61 alloy after isothermal annealing at 285 °C

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Fig. 11 TEM image of 5E61 alloy cold-rolled sheet after sensitization at 100 °C/168 h

at small angle grain boundaries, but the size of precipitation on small angle grain boundaries is limited, but the size of precipitates on large-angle grain boundaries is not suppressed, which increases the sensitivity at the large-angle grain boundaries. The precipitation of finely dispersed second phase particles in the matrix plays a significant pinning effect on dislocations, subgrain boundaries and grain boundaries, and can suppress dynamic recovery and recrystallization to reduce the number of large-angle grain boundaries, which is beneficial to improve the corrosion resistance of the alloy. Figure 11 is the TEM photograph of the 5E61 alloy cold-rolled sheet after sensitization at 100 °C/168 h. From Fig. 11a, a large number of deformed structures along the rolling direction can be observed. Moreover, the dislocation density in the alloy is high, and there are a large number of lath-like Al6 (Mn, Fe) and bean-shaped Al3 (Er, Zr) particles. In Fig. 11b, c, the β phase can be observed around the Al6 (Mn, Fe) phase and at the grain boundary, because the dislocation can reduce the precipitation activation energy of the grain boundary, and the second phase particles present in the matrix provide a low activation barrier for nucleation, providing a potential nucleation site for β phase formation. The formation of the primary β phase is generally promoted by dislocations to the periphery of the second phase particles and the grain boundaries. The initial β phase growth is expressed as isolated precipitates, but after a long period of sensitization annealing, it gradually grows to form a continuous strip shape, when the β phase in the alloy is continuously distributed around the second phase particles and at the grain boundary, the alloy is prone to intergranular corrosion. Therefore, the corrosion resistance of the 5E61 alloy in the cold-rolled state after sensitization at 100 °C/168 h decreased drastically. Figure 12 is the TEM photograph of the 5E61 alloy after annealing at 270 °C/4 h + 100 °C/168 h. It can be seen from Fig. 12a that the 5E61 alloy after stabilization and sensitization annealing has recovered and formed a large amount of cell structure, and deformation tissue and deformation recovery tissue were observed to alternate. In Fig. 12b, c, it can be seen that Al6 (Mn, Fe) and Al3 (Er, Zr) particles are dispersed, which can effectively pin dislocation and suppress dynamic recovery and recrystallization. At the same time, the dislocation density of the alloy after annealing is reduced, and the dislocation entanglement is transformed into sub-crystal to further reduce the deformation energy storage in the alloy, thereby reducing the dislocation

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Fig. 12 TEM image of 5E61 alloy after treatment at 270 °C/4 h + 100 °C/168 h

pipeline and effectively suppressing the diffusion of the β phase. In Fig. 12d, it is observed that there is no continuous β phase precipitated around the Al6 (Mn, Fe) particles and the fracture, and the corrosion resistance of the alloy is better.

4 Conclusions (1) The stabilized annealing process can improve the overall performance of the 5E61 cold-rolled sheet. Although the strength after annealing was lower than that of the cold-rolled sheet state, the long-term intergranular corrosion resistance was significantly better than that of the cold-rolled state. (2) There is a large amount of fibrous deformation structure inside the cold-rolled alloy, and the dislocation density was high. After the stabilized annealing, the interior of the alloy underwent recovery and recrystallization, the dislocation density is greatly reduced, and the dislocation channel of the β phase diffusion was reduced, thereby improving the corrosion resistance of the alloy. (3) The strength of the alloy after the higher stabilization annealing temperature was lower than that of the alloy after the low-temperature annealing; the lower stabilization annealing temperature required a relatively long annealing time to obtain high-strength corrosion resistance. The β phase preferentially precipitated

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on the small angle grain boundary. After the low-temperature annealing for a short time, only a small part of the recovery and recrystallization occurred inside the alloy, and a large number of small angle grain boundaries existed, which increased the corrosion sensitivity. Acknowledgements The authors are pleased to acknowledge the financial support received from the following projects (in no particular order). The National Key Research and Development Program of China (2016YFB0300804 and 2016YFB0300801), and the National Natural Science Fund for Innovative Research Groups (Grant No. 51621003). The Construction Project for National Engineering Laboratory for Industrial Big-data Application Technology(312000522303). National Natural Science Foundation of China (No. 51671005 and 51701006), Beijing Natural Science Foundation (2162006) and Program on Jiangsu Key Laboratory for Clad Materials (BM2014006).

References 1. R. Kaibyshev, F. Musin, D.R. Lesuer, T.G. Nieh, Superlastic behavior of an Al–Mg alloy at elevated temperatures. Mater. Sci. Eng. A 342(1–2), 169 (2003) 2. S. Katsas, J. Nikolaou, G. Papadimitriou, Corrosion resistance of repair welded naval aluminium alloys. Mater. Des. 28(3), 831 (2007) 3. Z.R. Nie, S.P. Wen, H. Huang et al., Research progress of niobium microalloyed aluminum alloy. Chin. J. Nonferr. Metals 21, 2361–2370 (2011) 4. G.F. Xu, J.J. Yang, T.N. Jin et al., Effect of trace rare earth Er on microstructure and properties of Al–5Mg alloy. Chin. J. Nonferr. Metals 21, 2361–2370 (2011) 5. L. Tan, T.R. Allen, Effect of thermo mechanical treatment on the corrosion of AA5083. Corros. Sci. 52(2), 548 (2010) 6. R.K. Gupta, R. Zhang, C.H.J. Davies et al., Influence of Mg content on the sensitization and corrosion of Al-xMg(-Mn) alloys. Corrosion 69(11), 1081–1087 (2013) 7. Y. Zhu, D.A. Cullen, S. Kar et al., Evaluation of Al3 Mg2 , precipitates and Mn-rich phase in aluminum-magnesium alloy based on scanning transmission electron microscopy imaging. Metall. Mater. Trans. A 43(13), 4933–4939 (2012) 8. D.S. D’Antuono, J. Gaies, W. Golumbfskie et al., Grain boundary misorientation dependence of β phase precipitation in an Al–Mg alloy. Scripta Mater. 76, 81–84 (2014)

Flow Stress Behavior and Microstructural Evolution of a High-Alloying Al–Zn–Mg–Cu Alloy Guohui Shi, Yong’an Zhang, Xiwu Li, Shuhui Huang, Zhihui Li, Lizhen Yan, Hongwei Yan and Hongwei Liu

Abstract The flow behavior of a high-alloying Al–Zn–Mg–Cu alloy was studied by compression tests with the temperature range of 300–440 °C and the strain rates range of 0.001–1 s−1 , and the corresponding microstructural evolution was observed. Results show flow stress curves exhibit the peak value at a critical strain, and the peak stress decreases with increasing of deformation temperatures. Numerous precipitated particles with a small size and high-density dislocations should be responsible for the high flow stress. Dynamic recovery is the main way of flow softening while dynamic coarsening of precipitated particles and dynamic recrystallization also play a role in flow softening under low-temperature and high-temperature conditions, respectively. The continuous dynamic recrystallization is the major mechanism for dynamic recrystallization behavior.

1 Introduction Al–Zn–Mg–Cu alloys have been widely used for aerospace structural components, such as wing ribs, spars, and fuselage, due to high strength, low density, better stress corrosion resistance and fracture toughness [1, 2]. To obtain desired microstructures and excellent properties, the Al–Zn–Mg–Cu alloy has to be subjected to a complex fabrication route including hot deformation and thermal treatment [3]. Hot deformation is a very complicated process for the interaction between force and heat during microstructure evolution [4]. Deformation temperature, deformation rate, and deformation degree are three basic parameters that control microstructural evolution [5–7]. Deformation temperature significantly affects dislocation slip and particle transformation, which are the major parts of the microstructural evolution due to the thermal activation of atomic migration. A great number of researches have been investigated the effect of deformation temperature on the microstructural evolution of the Al–Zn–Mg–Cu alloy during hot deformation. Deng [8] found the deformation G. Shi · Y. Zhang (B) · X. Li · S. Huang · Z. Li · L. Yan · H. Yan · H. Liu State Key Laboratory of Nonferrous Metals and Processes, GRINM Group Co., Ltd., Beijing 101047, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_11

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temperature has an evident influence on the size of subgrains and the distribution of dislocations increases with the increase of deformation temperature and the decrease of the strain rate, and revealed the dislocation characteristic. Feng [9] proposed that the softening mechanisms of homogenized Al–7.68Zn–2.12Mg–1.9Cu–0.12Zr alloy is a dynamic recovery and partial dynamic recrystallization at high temperature. Recently, for further strengthening the Al–Zn–Mg–Cu alloy, the content of the main alloying elements is increased further. Compared with the commonly commercial Al–Zn–Mg–Cu alloy, the studied alloy has higher Zn content (9.39 wt%) and Zn/Mg ratio, and lower Cu content. The chemical composition determines microstructure and properties of alloys, and the hot deformation microstructure is no exception. However, the hot deformation behavior of such high alloying Al–Zn–Mg–Cu alloys get little noticed by researchers, therefore, the study of the deformation behavior and the associated microstructure during compression at elevated temperature are necessary and contribute to understanding the deformation process of the high Zn-content aluminum alloy. Dynamic recovery (DRV) and dynamic recrystallization (DRX) are the typical softening mechanism in metals and alloys during hot deformation at elevated temperature [10]. For high stacking fault energy alloys, due to the high rate of DRV, the dynamic recrystallization is not easy to be observed. However, it was usually reported that DRX is discovered in aluminum alloys by many researchers [11–14], while two kinds of dynamic recrystallization mechanisms have been proposed generally: discontinuous dynamic recrystallization (DDRX) and continuous dynamic recrystallization (CDRX). The biggest difference between the two mechanisms is that whether there has a grain nucleation process or not. The DDRX involves the nucleation and growth of new grains with high angle boundaries. The CDRX refers to a way that subgrains with LAGB transform into the new grain with HAGBs [5, 13, 15–17]. Furthermore, the evolution of flow stress is correlated with different dynamic softening mechanisms during hot deformation at various deformation conditions [18]. In this work, the flow behavior and microstructural evolution of a high Zn content alloy were studied using uniaxial compression tests performed at various conditions. The interrelation between flow stress and microstructure were studied using electron backscattered diffraction (EBSD) technique and transmission electron microscopy (TEM). The microstructural evolution was investigated to reveal the softening mechanisms under various deformation conditions.

2 Experimental The experimental alloy was provided by Northeast Light Alloy Co. Ltd. in as-cast condition and its chemical composition is listed in Table 1. Before hot-compression tests, the alloy was homogenized at 470 °C for 24 h, and quenched to room temperature by water. The homogenized alloy was machined into cylindrical samples with 10 mm in diameter and 15 mm in height. Hot compression test was tested on the

Flow Stress Behavior and Microstructural Evolution … Table 1 The chemical composition of the high Zn content Al–Zn–Mg–Cu alloy (wt%)

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Composition

Zn

Mg

Cu

Zr

Fe

Si

Content

9.39

1.92

1.98

0.1

0.05

0.015

Gleeble-1500D thermo-mechanical simulation unit within the range of temperatures from 300 to 440 °C. Samples were heated to set predefined deformation temperature with a heating rate of 5 °C/s. To eliminate the thermal gradient, samples were held at the set temperature for 3 min. After hot compression, the samples were water quenched immediately. The total strain for the whole is 1. The deformed samples were sectioned parallel to the compression axis along centerlines to observe the microstructures by using JEOL JSM 7001F scanning electron microscope (SEM), attached with the electron backscattered diffraction function (EBSD), and JEOL JEM-2010 transmission electron microscopy (TEM). TEM samples were prepared by mechanical grinding to 50 μm and then twin-jet electropolishing in a solution of 25% HNO3 + 75% methanol cooled at −30 °C with a voltage of (10–15) V. EBSD samples were mechanical grounded and polished, and then electropolished in a solution of 90% Alcohol and 10% HClO4 within 5–7 s with a voltage of 20–25 V. EBSD analyses were performed using the OIM Analysis 5 software. Accounting for the Orientation map with boundaries graph, the blue line represents high angle grain boundary whose boundary angle is greater than 15 degrees, the red line and green line represent the low angle boundary, which is in the range of 2°–5° and of 5°–15°, respectively.

3 Results and Discussion 3.1 Flow Stress Behavior A series of flow stress–true strain curves are presented in Fig. 1. It can be found the flow stress rises rapidly in the initial strain and reaches the peak value at a strain of εp . After that, the flow stress has a little decrease, expressed by the σ , until the end of the deformation. The peak flow stress is very sensitive to deformation temperatures and strain rates. The εp becomes more and more smaller with the deformation temperature increase, suggesting that this alloy will have a more quick dynamic softening response under the higher temperature. The value of σ is calculated by (σ p − σ ε  0.5) and the figure of σ versus T is presented in Table 2. From which it can be found the value of σ decreases with the T increase. The lowest peak flow stress is 14 MPa at 440 °C and 0.001 s−1 and the highest one is 167 MPa at 300 °C and 10 s−1 . The peak flow stress decreases with the increase of deformation temperature, and it represents the deformation resistance of alloys under the specific deformation condition. From

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Fig. 1 Flow stress curves of the high Zn content Al–Zn–Mg–Cu alloy at different strain rates, a š  0.01 s−1 , b š  1 s−1 Table 2 The value of εp , σ p, and σ (Mpa) at various deformation temperature conditions

Temperature 300 °C

320 °C

360 °C

400 °C

440 °C

εp

0.0674

0.0613

0.0586

0.0344

0.0214

σ p (Mpa)

92.41

76.18

54.57

38.48

25.26

σ (Mpa)

5.69

5.38

4.49

1.27

1.44

the flow behavior analysis, it can be found the alloy has good workability and stability when the deformation temperature is high and the strain rate is low. Flow stress–true strain curve is a result of the competitive process of the dynamic softening and work hardening and reflects the microstructure evolution during hot deformation. Work hardening is related to dislocation increment with resulting in flow stress increase, while dynamic softening leads to flow stress decreases with respect to dislocation decrement. At the peak stress in εp , microstructural evolution attains a balance between work hardening and dynamic softening with a certain amount of dislocation. Elevated deformation temperature can reduce the value of critical resolved shear stress (CRSS), which weakens the work hardening effect and strength the dynamic softening. So, the higher deformation temperature, the lower the peak stress and the εp . After the εp , a little decrease of flow stress indicates dynamic softening is slightly stronger than the work hardening at this situation and the difference between the dynamic softening and work hardening is quantitatively present by σ . From the result, it is certain that the high deformation temperature can reduce the value of σ . Based on the above analysis, it can be concluded the dynamic softening really occurred and was sensitive to deformation temperature. To deeply understand the hot deformation under various temperatures, the microstructure characteristics and its evolution brought by dynamic softening are presented and discussed in the next section.

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Fig. 2 The microstructure of the homogenized high Zn content Al–Zn–Mg–Cu alloy: a the optical microstructure; b the figure of the orientation map

3.2 Microstructural Evolution Homogenization Microstructure Figure 2 shows the microstructure of the specimens homogenized at 470 °C for 24 h. The homogenized microstructure consists of equiaxial grains with an average diameter of about 117.3 μm and their boundaries were mostly characterized by high angle boundaries (HAGB). A few of second phases distribute at grain boundaries, most of which were identified as iron-rich phases by the energy Disperse Spectroscopy (EDS). No precipitated particles can be observed in the micrograph, indicating the Al matrix has a good solution of Zn, Mg, and Cu. Such an alloy with unstable microstructure often takes place dynamic precipitation and subsequent coarsening during hot deformation. Microstructural Evolution of Grains and Substructures Figure 3 shows the microstructure of specimens deformed under various deformation temperatures. The grains are compressed flat and were called as “elongated grains” also, whose boundaries are characterized by HAGB. Compared with initial grains, the interior of elongated grains was filled with the low angle grain boundaries (LAGB). The elongated grains are divided into several parts with different orientations by LAGB. Although dynamic recrystallization is rarely observed in metals with high stacking fault energy, there are a few recrystallized grains are found in the micrograph of the specimen deformed under high deformation temperature, as showed in Fig. 3d. The deformation temperature really controls the characteristic of those deformed microstructures. Table 3 lists the mean aspect ratio of elongated grains to estimate deformation degree of initial grains in plane paralleled to the axis of the cylinder specimen. It can be found that the low deformation temperature helps to reduce the vertical intercept of grains. From Fig. 3, it is easy to find that the morphology of LAGB is distinctly inconsistent at various deformation conditions. Under the high-deformation temperature condition, the distinct and straightness LAGB enclose into subgrains which are good at balancing of the mechanical property and corrosion resistance [8]. However,

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Fig. 3 Figures of orientation map for specimens deformed at various hot deformation temperatures with a strain rate of 0.01 s−1 , a 320 °C; b 360 °C; c 400 °C; d 440 °C Table 3 The mean aspect ratio and the diameter max. and diameter min. ratio of elongated grains

Temperature Aspect radio Dmax /Dmin

320 °C

360 °C

400 °C

440 °C

9.53

8.83

5.77

4.73

12.13

11.47

7.00

6.10

under the low deformation temperature, the LAGB are curly and mess, meanwhile, it is hard to distinguish whether there have subgrains are formed. For a better analysis of the microstructure characteristics, the EBSD technical provides quantitative measurements of grains orientation. Figure 4 is the statistical diagram of the LAGB and flow stress versus deformation temperatures. The fraction of LAGB decreases with the deformation temperatures increasing continuously, and the highest one is 69.6% with a temperature of 320 °C while the lowest one is 54.6% with a temperature of 440 °C. In addition, it can be seen the flow stress decrease with the deformation temperature increase. As is well-known, the boundary strength is a way to strengthen metal alloys, so the higher LAGB fraction, the higher efficiency of boundary strength, accompanied by higher flow stress. The LAGB fraction reflects the stored energy reserved in microstructure after hot deformation,

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Fig. 4 The flow stress and the LAGB fraction versus deformation temperatures, š  0.01 s−1

which affects the static recrystallization at the post-deformation treatment. There was another index which computes the average substructures orientation and judges the deformation of a grain: Grain Average Misorientation (GAM). Grains are classified as “deformed grains” when the GAM exceeds 1°, as “substructure grains” when the GAM is range from 0.5° to 1° and as “recrystallization grains” for the remaining grains whose GAM below 0.5°. The relative fractions of these three kinds of grains are shown in Fig. 5. The relative volume fraction of recrystallization grains seems has a valid value of 5%, only in the deformation with a temperature of 440 °C. The high deformation temperature also accelerates substructure grains formation. The relative volume fractions of substructure grains are 68 and 43% at deformation temperature of 440 and 400 °C, respectively, but only 5 and 3% at deformation temperature of 360 and 320 °C, respectively. Based on the above contrast in relative volume fractions of substructure grains and the observation on a micrograph of subgrains, it can be concluded that substructure grains are made up of subgrains which are formed through the dynamic recovery during hot deformation. The Deformed grain is filled up with curly and short LAGB and its relative volume fraction increase with the deformation temperature decrease. The LAGB is the primary product of dynamic recovery while the sub-grains are the ultimate production of the dynamic recovery. Based on the above results, it can be concluded dynamic recovery is more efficient at high deformation temperature conditions. The high deformation temperature gives high thermal energy so that the rate of dislocation slip and arrangement is high, resulting in larger and complete subgrains formed. For the mechanism of dynamic recrystallization of aluminum alloys, two kinds of mechanism can be recognized, one is continuous dynamic recrystallization (CDRX), another is discontinuous dynamic recrystallization (DDRX). In brief, CDRX refers to form recrystallization grains through subgrains transformation while DDRX refers to form recrystallization grains by grain nucleation. Obviously, the mechanism of DDRX involves high driving energy to form nucleation. Under a high-deformation temperature condition, the strain distortion energy of the corresponding microstructure is low. Therefore, it can be concluded the main DRX

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Fig. 5 The relative volume fraction of three microstructural structures deformed specimens

mechanism is CDRX. Large and complete subgrains are formed quickly by dynamic recovery, and then a few of them keep on rotating until recrystallization grains formation. However, due to the fraction of dynamic recrystallization grains is tiny under high deformation temperature conditions, its contribution to flow stress softening is limited. Therefore, the dynamic recovery is the main softening mechanism during hot deformation, dynamic recrystallization occurs only in the high deformation temperature range. Dislocations and Precipitated Particles Microstructure As shown in Fig. 6, they are bright-field TEM images of deformed samples. The substructures, including dislocations, precipitated particles and subgrains, can be observed in these images. It can be found deformation temperature has a huge influence on the number of dislocations and precipitated particles. At low deformation temperatures of 320 and 360 °C, the dislocations are dense and the precipitated particles are numerous and fine. At the high deformation temperatures of 400 and 440 °C, only a few large particles and dislocations are discovered. Due to the pinning effect of precipitated particles on dislocation slip, precipitated particles cause high-density dislocations and stored strain distortion energy area which is in favor of particle stimulated nucleation of recrystallization (PSN) during. However, it is hard to find the recrystallization grains formed around precipitated particles in TEM graphs, as shown in Fig. 6a and b, which indicates these particles do not act as sites for PSN. So these precipitated particles must be responsible for flow stress increase. The size and number of precipitated particles depend on the deformation temperature and strain. As shown in Fig. 6, small precipitated particles get dissolved, whereas large particles are retained at the high temperature and the higher the deformation temperature, the larger the size and the fewer the number of particles. Compared with the particles morphology of deformed microstructure when strain at 0.05 and 0.8, it can be found the later possess the larger size and the lesser number. Particles coarsening is another characteristic of microstructure evolution which can explain why the flow stress has a decrease σ after strain of εp . Elevated temperature will weaken the particles coarsening

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Fig. 6 TEM micrographs of specimens deformed at various deformation temperature, š  0.01 s−1 a 320 °C; b 360 °C; c 400 °C; d 440 °C

efficiency, so the value of σ decreases as deformation temperature increasing. As shown in Fig. 6c and d, the Al3 Zr particles can be found, but they are not tangled by dislocations. The Al3 Zr particles are thermostable and their size and distribution have few changes during hot deformation with various temperatures. Some researchers have demonstrated Al3 Zr particles effectively impede DRX during hot deformation, so the absence of recrystallization grains in some conditions maybe blame to Al3 Zr particles.

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4 Conclusion (1) The flow stress curves exhibited peak stress at critical strain, after which the flow stress decreased slightly until the end of deformation. The peak stress decreased with increasing of deformation temperatures and decreasing of strain rates. (2) The fraction of LAGB decreased as the deformation temperature is increasing. When the deformation temperature was high enough, large and complete subgrains could form in elongated grains. A few of recrystallization grains were found in the micrograph under deformation temperature of 440 °C. The precipitated particles became larger and lesser with increasing deformation temperature and stain. (3) The dynamic recovery was the main way for softening during hot deformation and its efficiency was controlled by the deformation temperature. Precipitated particles coarsening was one kind of microstructural evolution that contributed to flow softening, especially at low-deformation temperature conditions. Under the high-deformation temperature condition, dynamic recrystallization was responsible for flow softening also, however, its contribution to flow stress softening was limited due to the tiny fraction of dynamic recrystallization. Acknowledgements This study was financially supported by the National Key R&D Program of China (No. 2016YFB0300803, 2016YFB0300903), the National Program on Key Basic Research Project of China (No. 2012CB619504) and National Natural Science Foundation of China (No. 51274046).

References 1. H.J. Mcqueen, Development of dynamic recrystallization theory. Mater. Sci. Eng. A 387–389, 203–208 (2004) 2. T. Dursun, C. Soutis, Recent developments in advanced aircraft aluminium alloys. Mater. Des. (1980–2015) 56, 862–871 (2014) 3. M. Dixit, R.S. Mishra, K.K. Sankaran, Structure–property correlations in al 7050 and al 7055 high-strength aluminum alloys. Mater. Sci. Eng. A 478(1–2), 163–172 (2008) 4. H.J. Mc Queen, S. Spigarelli, M.E. Kassner, Hot Deformation and Processing of Aluminum Alloys (CRC Press, Boca Raton, 2016) 5. X. Liu, S. Han, L. Chen et al., Flow behavior and microstructural evolution of 7a85 highstrength aluminum alloy during hot deformation. Metall. Mater. Trans. A 48(5), 2336 (2017) 6. C.J. Shi, J. Lai, X.G. Chen, Microstructural evolution and dynamic softening mechanisms of al-zn-mg-cu alloy during hot compressive deformation. Materials 7(1), 244 (2014) 7. X.Y. Liu, Q.L. Pan, Y.B. He et al., Flow behavior and microstructural evolution of al–cu–mg–ag alloy during hot compression deformation. Mater. Sci. Eng. A 500(1–2), 150–154 (2009) 8. Y. Deng, Z. Yin, J. Huang, Hot deformation behavior and microstructural evolution of homogenized 7050 aluminum alloy during compression at elevated temperature. Mater. Sci. Eng. A 528(3), 1780–1786 (2011) 9. D. Feng, X.M. Zhang, S.D. Liu et al., Constitutive equation and hot deformation behavior of homogenized al–7.68zn–2.12mg–1.98cu–0.12zr alloy during compression at elevated temperature. Mater. Sci. Eng. A. 608, 63–72 (2014)

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10. H.J. Mc Queen, J.J. Jonas, Recent advances in hot working: fundamental dynamic softening mechanisms. 3, 233–241 (1984) 11. Z. Fei, S. Jian, Y. Xiaodong et al., Constitutive analysis to predict high-temperature flow stress in 2099 al-li alloy. Rare Metal Mat. Eng. 43(6), 1312–1318 (2014) 12. Q. Yang, X. Wang, X. Li et al., Hot deformation behavior and microstructure of aa2195 alloy under plane strain compression. Mater. Charact. 131, 500 (2017) 13. M. Zhou, Y.C. Lin, J. Deng et al., Hot tensile deformation behaviors and constitutive model of an al–zn–mg–cu alloy. Mater. Des. 59, 141 (2014) 14. Y. Wei, B.Q. Xiong, Y.A. Zhang et al., Research on flow stress of spray formed 70si30al alloy under hot compression deformation. Rare Met. 6(25), 665 (2006) 15. C. Hi, X.G. Chen, Evolution of activation energies for hot deformation of 7150 aluminum alloys with various zr and v additions. Mater. Sci. Eng. A. 650, 197–209 (2016) 16. J. Li, J. Shen, X. Yan et al., Microstructure evolution of 7050 aluminum alloy during hot deformation. T Nonferr Metal Soc. 20(2), 189–194 (2010) 17. N. Jin, H. Zhang, Y. Han et al., Hot deformation behavior of 7150 aluminum alloy during compression at elevated temperature. Mater. Charact. 60(6), 530 (2009) 18. C. Huang, J. Deng, S.X. Wang et al., A physical-based constitutive model to describe the strainhardening and dynamic recovery behaviors of 5754 aluminum alloy. Mater. Sci. Eng. A 699, 106 (2017)

Influence of Space Environment on the Properties of Diamond/Cu Composites Zhongnan Xie, Hong Guo and Ximin Zhang

Abstract In this paper, several space environmental simulation tests, including particle irradiation and thermal cycling were carried out on the Diamond/Cu composites. The properties and microstructure of composites before and after different environmental effects were compared. The results show that the thermal conductivity of the composites is above 600 W/mK and the change amplitude is less than 5% after the space environment simulation tests. The bending strength is higher than 400 MPa after space environmental simulation tests, and the bending strength is increased by about 15% after particle irradiation. After the space environment test, the Diamond/Cu composite still maintains good thermal physical properties and mechanical properties, which meet the requirements of the space environment.

1 Introduction Diamond/Cu composites are expected to become the new-generation thermal management materials for aerospace electronic devices, due to their lower specific gravity, excellent thermal conductivity, and adjustable thermal expansion coefficient [1, 2]. In recent years, Diamond/Cu composites have attracted extensive attention owing to their excellent thermal conductivity. At present, the thermal conductivity of Diamond/Cu composites at product level has reached 600 W/mK [3, 4]. Compared with the traditional thermal management materials, such as SiCp /Al composites, tungsten copper, and molybdenum copper, the thermal conductivity of Diamond/Cu composites has been increased by 2–3 times [5–7]. Therefore, the application of Diamond/Cu composites in space electronic devices can effectively improve the heat dissipation of its core components and prolong the life of space electronic devices. Z. Xie · H. Guo (B) · X. Zhang General Research Institute for Nonferrous Metals, Beijing 100088, China e-mail: [email protected] Z. Xie · H. Guo · X. Zhang National Engineering Research Center for Nonferrous Metals Composites, Beijing 100088, China © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_12

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At present, there is no report on the performance changes of Diamond/Cu composites for aerospace electronic devices after undergoing space environmental testing. Therefore, in view of the application conditions of the thermal management materials for space electronic devices, the changes in the physical and mechanical properties of Diamond/Cu composites after the ground simulation tests (including thermal cycling and particle irradiation tests) were studied [8, 9]. Combining with the structural characteristics of Diamond/Cu composite, the law of performance change before and after tests is analyzed.

2 Experimental Procedure The preparation method of Diamond/Cu composites used in the experiment was reported in Ref. [10]. The composite material of diamond volume fraction of 60 vol.% was prepared by pressure infiltration method, and then the samples for thermal diffusivity and bending strength were manufactured by mechanical processing (the size and quantity were carried out according to the related test standards). The above Diamond/Cu composite specimens were randomly divided into 3 groups, and the classification is shown in Table 1. First groups were taken as blank, second groups were tested with thermal cycling, the third groups were particle irradiation test. The thermal cycling test was carried out in accordance with standard GJB2502.8-2006. It was recycled 50 times in the range of −50 to 150 °C, the temperature changing rate was (8–10) °C/min, and the residence time was 15 min. The particle irradiation test was carried out in accordance with the standard QJ100042008. The radiation source was 60Co gamma ray with a total radiation dose of 2.3 × 105 rad (Si). The thermal physical and mechanical properties of the blank samples and the samples after the environment tests were tested respectively. The effects of different environmental effects on the properties of Diamond/Cu composites were analyzed. The surface morphology of Diamond/Cu composite was observed by Hitachi 5-4800 cold field emission scanning electron microscopy (SEM), and the interfacial bonding between diamond and copper matrix alloy was analyzed. The three-point bending strength of the prepared Diamond/Cu composites was measured by a universal material testing machine, with a span of 25 mm and a loading rate of 0.5 mm/min. The thermal conductivity of the material at room temperature was measured by laser flash method with the LFA447 thermal conductivity tester of NETZSCH Company.

Table 1 Classification table for environmental test samples

Groups

First

Second

Third

Test project

Nothing

Thermal cycling

Particle irradiation

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3 Results and Discussion 3.1 Microstructure of Fracture of Diamond/Cu Composites Figure 1 shows the fracture morphology of the Diamond/Cu composites including the blank group and the groups after different environmental tests. There are three main forms of flexural fracture of Diamond/Cu composites: the transgranular fracture of diamond particles, the toughened fracture of the matrix and the disintegration of the interface between the diamond and the matrix [11]. The three forms of fracture correspond to I, II, and III in the SEM images, as shown in Fig. 1a and c, respectively. When the interface between diamond and substrate is destroyed, the bonding strength between substrate and diamond is weak. When the diamond transgranular fracture or matrix toughening fracture occurs, the interface bonding strength is higher [12, 13].

Fig. 1 SEM images of sample fracture after different environmental tests a blank; b particle irradiation; c thermal cycling

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Fig. 2 Properties of Diamond/Cu composites after particle irradiation test

3.2 Effect of Particle Irradiation on Properties of Diamond/Cu Composites The properties of Diamond/Cu composites after particle irradiation test are shown in Fig. 2. Thermal conductivity of Diamond/Cu composites decreased by about 30 W/mK after particle irradiation test, which was reduced by 4.8%. The flexural strength increased from 418.5 to 484.9 MPa, which improved greatly. According to the bending fractured SEM images of samples before and after irradiation of particles in Fig. 1a and b, it is known that the main fraction of the samples is toughened fracture of the matrix, and there is a small amount of diamond extraction after the particle irradiation. It is shown that the interfacial bonding strength of diamond and matrix decreases after irradiation, and the interfacial thermal resistance increases with the weakening of the interface, so the thermal conductivity of the composite begins to decrease [14, 15]. Compared to the brittle fracture mode of the diamond transgranular fracture, the toughened fracture of the matrix and the extraction of the diamond particles need to overcome the higher external force. At the same time, the path of dislocation slip is increased, and more energy is consumed when the material is destroyed [8]. Therefore, the flexural strength of the composite is improved.

3.3 Effect of Thermal Cycling on Properties of Diamond/Cu Composites The performance of Diamond/Cu composite after high- and low-temperature cycling test is shown in Fig. 3. Thermal conductivity and bending strength of Diamond/Cu composites decreased after 50 times thermal cycling at −50 to 150 °C, and the decrease was 2.7 and 1.5%, respectively. The variation of the properties of the Diamond/Cu composites is less than 3%, indicating that the composites have strong resistance to alternating temperature.

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Fig. 3 Properties of Diamond/Cu composites after thermal cycling test

It is known from Fig. 1 that there is an obvious inhomogeneity in the microstructure of Diamond/Cu composites. There is a two-phase interface between the matrix and the reinforcement. The bonding strength of the interface, the composition, and thickness of the interface layer determine the performance of the Diamond/Cu composite. As shown in Fig. 1a and c, there is a more diamond transgranular fracture in the bending fracture in the blank group, while the bending fracture of the sample after the thermal cycling is mainly toughened fracture of the matrix, and there is a small amount of diamond transgranular fracture at the same time. It is shown that after the thermal cycling, the diamond and the matrix still have a high interfacial bonding strength, which shows a slight decrease in the thermal conductivity and the bending strength.

4 Conclusions (1) After the particle irradiation test, the thermal conductivity of Diamond/Cu composites decreased by about 30 W/mK and the decrement was 4.8%. The flexural strength increased from 418.5 to 484.9 MPa, which was improved by about 15%. (2) After 50 times high and low thermal cycling tests, the thermal conductivity and flexural strength of Diamond/Cu composites decreased slightly, and the decline rates were 2.7 and 1.5%, respectively. The change of performance was within 3%, indicating that the composite has a strong ability to resist the temperature variation. (3) After the space environment tests, the thermal conductivity of the composites was above 600 W/mK and the bending strength was above 400 MPa. It had a good ability to resist the change of the space environments and meets the requirements of the space environments.

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Acknowledgements The authors would like to acknowledge the financial support from the Ministry of Science & Technology of China (the National Key Research and Development Program of China No. 2016YFB0301402).

References 1. S.A. Coogan, J. Maynard, M. Stoklosa et al., in Thermal Management of a High Power Multichip Module in a Space Environment. Aip Conference (American Institute of Physics, 1996), pp. 949–954 2. S.V. Garimella, V. Singhal, D. Liu, On-chip thermal management with microchannel heat sinks and integrated micropumps. Proc. IEEE 94(8), 1534–1548 (2006) 3. Y. Zhang, H.L. Zhang, J.H. Wu et al., Enhanced thermal conductivity in copper matrix composites reinforced with titanium-coated diamond particles. Scripta Mater. 65(12), 1097–1100 (2011) 4. A.M. Abyzov, S.V. Kidalov, F.M. Shakhov, High thermal conductivity composite of diamond particles with tungsten coating in a copper matrix for heat sink application. Appl. Therm. Eng. 48(15), 72–80 (2012) 5. Y. Cui, L. Wang, J. Ren, Multi-functional SiC/Al composites for aerospace applications. 21(6), 578–584 (2008) 6. J. Fan, Properties of SiC_p/Al composites for aerospace application. Aerosp. Mater. Technol. (2005) 7. L. Zhang, W. Chen, G. Luo et al., Low-temperature densification and excellent thermal properties of W-Cu thermal-management composites prepared from copper-coated tungsten powders. J. Alloy. Compd. 588(588), 49–52 (2014) 8. Y. Nakayama, K. Imagawa, M. Tagashira et al., Evaluation and analysis of thermal control materials under ground simulation test for space environment effects. High Perform. Polym. 13(3), S433–S451 (2001) 9. T. Paulmier, B. Dirassen, D. Payan et al., Material charging in space environment: experimental test simulation and induced conductive mechanisms. IEEE Trans. Dielectr. Electr. Insul. 16(3), 682–688 (2009) 10. X. Zhang, H. Guo, F. Yin et al., Interfacial microstructure and properties of diamond/Cu-Cr composites for electronic packaging applications. Rare Met. 30(1), 94–98 (2011) 11. F. Yin, G. Hong, X. Zhang et al., Effect of Ti on microstructures and properties of Diamond/Copper composites. Acta Mater. Compos. Sinica 27(3), 138–143 (2010) 12. K. Chu, C. Jia, H. Guo et al., On the thermal conductivity of Cu–Zr/Diamond composites. Mater. Des. 45(45), 36–42 (2013) 13. P.P. Wang, H. Guo, X.M. Zhang et al., Research Progress of Interface on Diamond/Copper Composites for Thermal Management (2014), pp. 680–688 14. M. Zain-Ul-Abdein, Numerical investigation of the effect of interfacial thermal resistance upon the thermal conductivity of Copper/Diamond composites. Mater. Des. 86, 248–258 (2015) 15. H. Xin, A. Mosallam, Y. Liu et al., Impact of hydrothermal aging on rotational behavior of webflange junctions of structural pultruded composite members for bridge applications. Compos. B 110, 279–297 (2017)

The Effects of Electron Beam Welding Parameters on Microstructure and Properties of GH4738 Alloy Xinxu Li, Yong Zhang, Peihuan Li and Shaomin Lv

Abstract The GH4738 alloy is widely applied in different industries due to excellent strength and toughness matching stability and a low crack propagation rate. There are several welding methods used in GH4738 alloy. In order to increase the demand for high speed and low distortion welding, electron beam welding is introduced into jointing engineering pieces. The purpose of this study was to evaluate the effect of electron beam welding parameters (welding speed and beam current) on macroscopic morphology, the room temperature tensile properties, microstructure and microhardness of GH4738 alloy. According to the macroscopic morphology, there is a very narrow heat-affected zone during electron beam welding process. By analyzing microstructure, it could be concluded that the great temperature gradient in the small weld zone provides favorable conditions for the growth of the columnar crystal. Furthermore, the value of microhardness in the weld area is far lower than the value of microhardness in the base metal, especially in heat-affected zone. Eventually, according to these results, optimum electron beam welding parameters is selected. When the welding speed is 20 mm/s and beam current is 33 mA, there is more compact macroscopic morphology and better mechanical properties in the

X. Li (B) · Y. Zhang · P. Li · S. Lv Science and Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, Beijing 100095, China e-mail: [email protected] Y. Zhang e-mail: [email protected] P. Li e-mail: [email protected] S. Lv e-mail: [email protected] X. Li School of Metallurgy, Northeastern University, Shenyang 110819, China S. Lv Institute for Advanced Materials and Technology, University of Science and Technology Beijing, Beijing 100083, China © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_13

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weldment and the hot affecting zone is narrow. What’s more, the microhardness and the room tensile strength of weldments are excellent after heat treatment.

1 Introduction The GH4738 alloy (known as Waspaloy in the United States) is a superalloy developed via vacuum smelting by Special Metal Company in 1952 in the New Hartford for the first time. It is mainly equipped with the PWA Pratt & Whitney aero engine turbine blades, and then is selected as the engine turbine disc for the Boeing 727 and Boeing 747 aircraft [1]. It could be regarded as a traditional precipitation-strengthened superalloy with very good strength at high temperature [2]. Its high-temperature strength is derived mainly from the precipitates of the Ni3 (Al, Ti) phase [3]. Compared to GH4169, GH4738 improves the mass percent of Al and Ti, and reduces the content of Nb, which improves the comprehensive properties [4]. Furthermore, solid solution strengthening elements, Mo, Co, and Cr, have a minor high-temperature strength effect. The addition of carbon accelerates the formation of carbides in the form of MC and M23 C6 , which further increase the creep resistance through pinning the grain boundaries and reducing grain boundary sliding at high temperature [5, 6]. The GH4738 has moderate strength at low temperature and excellent high-temperature creep strength [7]. It has excellent strength and toughness matching stability, a low crack propagation rate and a good oxidation resistance [8]. Therefore, it has been widely applied in different industries such as the turbine disc, blades, bolts and other components of the gas turbine in aerospace and other fields [9]. Welding is the most efficient means of producing permanent joints. In order to increase demand for high speed and low distortion welding, electron beam welding has been introduced into jointing engineering pieces as a useful method [10]. Although competitive laser welding has been developing quickly, electron beam welding remains indispensable in many fields due to the possibility of obtaining greater penetration depth and its greater welding rate [11]. Due to the high vacuum environment inside the chamber, the material could be protected from contamination during the welding process [12, 13]. The welding method doesn’t have only small heat-affected zone and high efficiency, but also lacks the defects in the melting region and HAZ [14, 15]. Therefore, it is widely applied in automotive, nuclear, electrical engineering, aerospace and mechanical engineering [16]. At present, some researchers have deeply studied the electron beam welding process. Budkin studied some metals’ weldability under different welding conditions and developed the physical–mathematical model [17, 18]. Demyanov and Kaniukov found the superconducting properties of the welded specimens due to an increase in the concentration of impurities [15]. Although GH4738 superalloy has many important applications in advanced industries such as aerospace and military, very few researches and reports on the electron beam welding have been conducted so far. Therefore, it is very meaningful and necessary to do some researches about electron beam welding processing for GH4738

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alloy. In this study, the effect of electron beam welding parameters on microstructure, weld penetration depth and mechanical properties of GH4738 alloy would be investigated and discussed further. Moreover, the weldability of metal materials is a little limited by susceptibility to post-welding heat treatment cracking [19]. Accordingly, the effect of heat treating on microhardness of GH4738 alloy would also be studied.

2 Experimental Procedures The GH4738 alloy was melted through vacuum induction and vacuum arc remelting, and the ingot was forged, rolled, and machined. The test alloy was a hot-rolled plate through solid solution treatment, and the plate was 10 mm thickness. The chemical composition of the material is shown in Table 1. In order to remove the dirt of hot rolled plate surface, the material used should be carefully and strictly cleaned by sandpaper and then they would be immersed into acetone solution for 6 s. Electron beam welding was performed at a beam voltage of 150 kV, focus current of 2200 mA, different welding speed and different beam current in GH4738 hot-rolled plate, as shown in Table 2. After welding, the weldments were undergoing heat treating with the regime of 845 °C × 4 h, AC + 769 °C × 16 h, AC. In order to study and compare the properties of electron beam welded joint with the base metal of GH4738 alloy, the tensile samples were prepared with the fusion zone parallel to the loading direction. The specification of tension samples was φ5 * 25 mm

Table 1 Chemical composition (mass %) of GH4738 C

Si

Mn

Cr

Co

Mo

Al

Ti

Zr

Ni

0.06

0.12

0.79

19.54

13.55

4.45

1.37

2.89

0.075

Bal.

Table 2 The parameters of BEW processing Serial number

Voltage U B (kV)

Focus current I L (mA)

Welding speed v (mm/s)

Beam current I B (mA)

1

150.0

2200

30

40 44 36

2

150.0

2200

20

30 33 27

3

150.0

2200

10

20 22 18

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and K  5.65. Microhardness measurements were conducted at a regular spacing (300 μm) across the weldment from base metal on one side to base metal on the other side. To study the effect of electron beam welding parameters on the microstructure of different welding regions, the welded samples were cut from the cross-section containing all of the weld regions. The samples were prepared in a metallographic procedure. Samples were etched with a special solution. The samples were first examined for the geometry depth of penetration of the welds and microstructure by an optical microscope (OM).

3 Results The Macroscopic Morphology of Weldments The macroscopic morphology of the cross-section of weldments was investigated under different welding speed and beam current, as presented in Fig. 1. According to the macroscopic morphology, the feature of the narrow heat-affected zone could be significantly observed after the electron beam welding process. What’s more, it could be observed that when the welding speed was 30 mm/s, there were obvious cracks in the weldments. However, no crack appeared in the weldments when the welding speed slowed down below 20 mm/s. How to explain this phenomenon? Electron beam welding is a special metallurgical. While the melting metal starts to cool from a liquid into solid, the inner stress increases during this process [20]. Faster welding speed is equivalent to shorter melting time. In shorter melting time, inner stress increase quickly. So there were obvious cracks when the welding speed was 30 mm/s. At the same time, when the welding speed was 10 mm/s, a significantly wider welding zone is formed than the width at the speed of 20 and 30 mm/s, because the heat emission speed is relatively slow at the low welding speed. According to the welding depth and welding joint morphology, the optimal beam current would be chosen. Two conditions must be satisfied in the process of electron beam welding: (1) Suitable welding depth. Electron beam welding is a high-energy welding process with high energy density and high welding depth. However, if the weld depth is too large and extends to the pioneer wire area, the pioneer wire under the welding mating surface would be seriously damaged, resulting in poor performance of the composite material. So it is necessary to ensure the suitable welding depth. (2) Ensure that the welded joint is compact and free of pores. It is important to ensure that welded joint is close-grained and there is no existence of the pore clearance. If the leakage situation occurs, the hot isotactic pressing process would not be able to generate the effect of densification. This process is equivalent to heat treatment process under inert gas protection, whose influence on the sample performance is disastrous. It could be seen from Fig. 1, the depth of the weld layer varied along with the change of beam current obviously. According to the two conditions above, the

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Fig. 1 Macroscopic morphology of welded joints with different welding parameters: a welding speed of 30 mm/s; b welding speed of 20 mm/s; c welding speed of 10 mm/s

corresponding optimum beam current 44, 33, and 22 mA could be chosen at the welding speed 30, 20 and 10 mm/s, respectively. Tensile Properties of Weldments In order to investigate the strength of the weld joints, the room temperature tensile properties of GH4738 weldments were tested under different welding speed (Fig. 2). The tensile results include the data before heat treatment and the data after heat treatment. It can be seen that there are more excellent mechanical properties under the welding speed of 20 mm/s. The result fits neatly into the macroscopic morphology. At the same time, it can be observed that heat treatment plays a very significant role in improving the weldments’ strength. It could be seen from Fig. 2, the tensile strength of the weldments is excellent after heat treatment, more than 85% of the base metal strength. The inner grain organization would be investigated further in the passage below. Analysis of Weldments Microstructure Taking a sample at the welding speed v  20 mm/s and the beam current I B  30 mA to analyze microstructure evolution after electron beam welding, as shown in Fig. 3. It can be observed that the weld shape of weldment is an axis-symmetric nails shape. The size of the head area is about 4 mm in height and about 2 mm in depth. The middle of the weldment is uniform and the weld width is about 1.25 mm. The ratio of depth and width can reach more than 5–1. Lower weld region presented a slender spike shape due to the temperature field characteristics in the process of welding. Figure 3c, d shows the microstructure morphology of the weld zone, and it can be seen that the weld zone formed a typical dendritic structure. Besides, the growth of solidification grains in the weld area is perpendicular to the fusion boundary due to heat transfer in a signal

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Fig. 2 Tensile strength at room temperature of weldments with different parameters before and after heat treatment

direction. Figure 3b shows the microstructure picture of the weld joint area between the weld zone and the base material. The organization of the base metal and the weld metal combined very well. There was no obvious boundary and also no defect between the base metal and the weld metal. Meanwhile, no significant change is observed about the size of the base metal grain because not much heat is transferred to the base metal zone around the weld metal. The solidification mode of the electron beam welding changes along with the variety of welding technology process. The solidification mode is mainly decided by the growth rate and temperature gradient G/R. With the G/R ratio gradually reduce, solidification way in the fusion zone change from the flat growth into cellular growth, the growth of the columnar crystal, until the equiaxial crystal growth. Electron beam welding process belongs to the high quantity welding way, whose temperature of the center molten pool can reach 1800–2300 °C. High temperature in short time and fast cooling speed prompted the small welding area. The great temperature gradient in the small weld zone provides favorable conditions for the growth of the columnar crystal. So it could be observed in Fig. 3c, d that typical characteristics of columnar dendrites formed. From the weld metal zone to the base metal zone, the feature of heat-affected zone is not very obvious because the heat is lost faster during the welding process and the grain in the weld area close to base metal is no time to grow up. The grain size of the base metal area remained the same with the original parent metal. Analysis of Weldments Microhardness Figure 4 presents the profile of microhardness variation. It can be seen that the value of microhardness in the weld area is far lower than that in the base metal. This result is consistent with the conclusion of Ali [21]. Heat-affected zone between the weld and parent metal thickness is about 0.5 mm, where microhardness value is the lowest. The sharp drop in the heat-affected zone caused by the increase of the temperature is due to the increase in the grain size

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Fig. 3 The microstructures of the interface layer welded at a welding speed of 20 mm/min and beam current of 30 mA Fig. 4 The microhardness of welded joints before and after heat treatment

[22]. The phenomenon can also be observed from Fig. 3. Furthermore, there is an obvious increase in microhardness after heat treatment. It is a very promising prospect that the electron beam welding applied in the GH4738 alloy. According to the investigation above, the advantage of electron beam welding would be summarized. According to the macroscopic morphology analysis,

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very narrow heat-affected zone could be significantly seen after electron beam welding process. Compact and no crack organization would appear and appropriate welding depth and width in the weldments would be formed in proper electron beam welding parameters. The result of the room temperature tensile properties GH4738 weldments fitted neatly into the macroscopic morphology. The organization of the base metal and the weld metal combined very well. There was no obvious boundary and also no defect between the base metal and the weld metal. Meanwhile, no significant change was observed about the size of the base metal grain because not much heat was transferred to the base metal zone around the weld metal. Heat-affected zone was very narrow because the heat was lost faster during the welding process. Furthermore, there were more excellent properties after heat treatment, the room temperature tensile and microhardness.

4 Conclusion In this study, the influence of different electron beam welding speed and beam current on macroscopic morphology, the room temperature tensile properties, microstructure and microhardness of GH4738 alloy were investigated. The main results of this survey are concisely summarized as follows: (1) When the welding speed was 30 mm/s, there were obvious cracks in the weldments. However, no crack appeared in the weldments when the welding speed was slowed down below 20 mm/s. A significantly wider welding zone was formed when the welding speed was 10 mm/s. (2) The sample presented excellent mechanical properties at 20 mm/s. After heat treatment, the tensile strength of the weldments was more than 85% of the base metal strength. (3) Short time at high temperature and fast cooling speed prompted the small welding area. The great temperature gradient in the small weld zone provided favorable conditions for the growth of the columnar crystal. (4) There was a lower microhardness value in the weld area than in the base metal. Heat-affected zone between the weld and parent metal thickness was about 0.5 mm, where microhardness value was the lowest. Furthermore, there was an obvious increase in microhardness after heat treatment.

References 1. Z. Yao, J. Dong, M. Zhang, Gamma prime phase evolution during long-time exposure for GH738 superalloy. J. Trans. Mater. Heat Treat. 34, 31–37 (2013) 2. Z. Yao, M. Zhang, J. Dong, Stress rupture fracture model and microstructure evolution for Waspaloy. J. Metall. Mater. Trans. A 44, 3084–3089 (2013)

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3. A. Amiri, S. Bruschi, M.H. Sadeghi et al., Investigation on hot deformation behavior of Waspaloy. J. Mater. Sci. Eng. A 562, 77–82 (2013) 4. Y. Hu, D. Liu, X. Zhu et al., Effect of rolling passes on thermal parameters and microstructure evolution via ring-rolling process of GH4738 superalloy. J. Int J Adv Manuf. Technol. 96, 1165–1174 (2018) 5. C.G. Roberts, S.L. Semiatin, A.D. Rollett, Particle-associated misorientation distribution in a nickel-base superalloy. J. Scripta Mater. 56, 899–902 (2007) 6. A. Chamanfar, M. Jahazi, J. Gholipour et al., Evolution of flow stress and microstructure during isothermal compression of Waspaloy. J. Mater. Sci. Eng. A 615, 497–510 (2014) 7. K.M. Chang, X. Liu, Effect of γ’ content on the mechanical behavior of the WASPALOY alloy system. J. Mater. Sci. Eng. 735, 928–937 (2001) 8. Z. Yao, J. Dong, M. Zhang et al., Hot deformation behavior of superalloy GH738. J. Rare Metal Mater. Eng. 42, 1199–1204 (2013) 9. Z. Yao, J. Dong, M. Zhang, Microstructure control and prediction of GH738 superalloy during hot deformation I. Construction of microstructure evolution model. J. Acta Metall. Sinica 47, 1581–1590 (2011) 10. B.S. Yilbas, M. Sami, J. Nickel, Introduction into the electron beam welding of austenitic 321-type stainless steel. J. Mater. Process. Technol. 82, 13–20 (1998) 11. M.St. W˛eglowski, S. Błacha, A. Phillips, Electron beam welding-techniques and trends-Review. J. Vac. 130, 72–92 (2016) 12. H. Schultz, Electron Beam Welding. Cambridge (1993) 13. C.Y. Ho, Fusion zone during focused electron beam welding. J. Mater. Process. Technol. 167, 265–272 (2005) 14. Y. Wang, P. Fu, Y. Guan, Research on modeling of heat source for electron beam welding fusion-solidification zone. J. Chin. J. Aeronaut. 26, 217–223 (2013) 15. E. Demyanov, E.Y. Kaniukov, Superconducting properties of ultra-pure niobium welded joints. J. Low Temp. Phys. 41, 522–527 (2015) 16. X.Z. Li, S.B. Hu, J.Z. Xiao et al., Effects of the heterogeneity in the electron beam welded joint on fatigue crack growth in Ti-6Al-4V alloy. J. Mater. Sci. Eng. A 529, 170–176 (2011) 17. Y.V. Budkin, Special features of the physical-chemical processes of interaction of refractory metals in electron beam welding. J. Weld. Int. 25, 309–312 (2011) 18. Y.V. Budkin, V.A. Erofeev, Physic-mathematical model of the process of electron beam welding refractory metal to steel. J. Weld. Int. 25, 562–565 (2011) 19. Z. Li, S.L. Gobbi, J.H. Loreau, Laser welding of Waspaloy sheets for aero-engines. J. Mater. Process. Technol. 65, 183–190 (1997) 20. F. Gao, P. Li, P. Jiang et al., The effect of constraint on microstructure and properties of titanium alloy electron beam welding. J. Mater. Sci. Eng. A 721, 117–124 (2018) 21. H. Ali, N.-M. Homam, Effect of beam welding current variations on the microstructure and mechanical properties of Nb-1Zr advanced alloy. J. Vac. 150, 196–202 (2018) 22. H. Ali, N.-M. Homam, Electron beam welding of difficult-to-weld austenitic stainless steel/Nbbased alloy dissimilar joints without interlayer. J. Vac. 2017, 170–178 (2017)

Microstructure and Grain Refining Performance of High-Quality Al–5Ti–1B Master Alloy Yuehua Kang, Shuncheng Wang, Nan Zhou and Dongfu Song

Abstract The chemical composition, microstructure, and grain refining performance of high-quality Al–5Ti–1B master alloy were investigated. The results show that the content of alloying elements Ti and B in the Al–5Ti–1B master alloy is stable. The content of impurity elements Fe, Si, V, and K is very low. The microstructure of Al–5Ti–1B master alloy is uniform, fine and no oxide inclusions. The average size of TiAl3 and TiB2 phases is 16.7 and 0.73 μm, respectively. By adding 0.2 wt% Al–5Ti–1B master alloy, the grain size of pure Al has been refined to 75.7 μm, showing an excellent grain refining efficiency, strong ability to anti-fading and good adaption to a wide range of aluminum-melt temperature.

1 Introduction In the industry producing aluminum and its alloys, it is often first critical to obtain a fine equiaxed grain structure by adding grain refiner [1, 2]. A variety of benefits could be achieved such as an improvement of mechanical properties, formability and machinability, and so on. The Al–5Ti–1B master alloy becomes the most commonly used grain refiner. As reported that 75% of the product of aluminum and its alloys has been added the grain refiner of Al–5Ti–1B master alloy [3]. Various quality grades of Al–5Ti–1B master alloy are selected to produce different aluminum products. As to the high-precision aluminum product such as foil, cans and PS version, the

Y. Kang · S. Wang (B) · N. Zhou · D. Song Guangdong Institute of Materials and Processing, Guangzhou 510650, China e-mail: [email protected] Y. Kang e-mail: [email protected] N. Zhou e-mail: [email protected] D. Song e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_14

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Al–5Ti–1B master alloy needs a higher requirement of metallurgical quality and grain refining performance. However, the production of high-quality Al–5Ti–1B master alloy rods is difficult. Traditionally, the high-quality Al–5Ti–1B master alloy rods are mainly imported from abroad [4]. In recent years, we have successively developed multilayer multifrequency coil electromagnetic induction melting technology, purification technology of alkaline earth metal fluoride melt and high-strain high-speed continuous casting rolling technology [5–7]. The quality of the Al–5Ti–1B master alloy rods produced by us has reached the international leading level. Thus, in this paper, the chemical composition, microstructure, and grain refining performance of the high-quality Al–5Ti–1B master alloy rods produced by us has been investigated. It could be meaningful for grain refiner manufacturers and aluminum processing enterprises in selecting and using high-quality Al–5Ti–1B master alloy rods.

2 Experimental Procedure Three batches of Al–5Ti–1B master alloy rods with a diameter of 9.5 mm produced by our research group were investigated. The chemical composition of the alloy rods were analyzed by JY-ULTIMA2 inductively coupled plasma atomic emission spectroscopy (ICP-AES), and the results are listed in Table 1. The microstructures were observed by Leica DMI3000M optical microscopy (OM) and the morphology and size of the TiAl3 and TiB2 particles in the microstructure was further observed and measured by JCXA-733 scanning electron microscopy (SEM). Industrial pure aluminum (99.7 wt%) was used to perform the grain refining performance of the Al–5Ti–1B master alloy rods. It was melted in a well-type resistance furnace with a graphite crucible and cast into a cylindrical steel mold (outside diameter × height × thickness: 75 mm × 25 mm × 5 mm) with a silicon dioxide foam [8]. Three experiment groups were designed. The first group: four different contents of 0.05, 0.1, 0.2, and 0.3 wt% Al–5Ti–1B master alloy rods were added into the aluminum melt when the temperature was 720 °C, respectively. After 10 min of stirring and holding, the melt was cast into the mold. The second group: six different holding times of 1, 2, 10, 30, 60, and 120 min were processed after addition

Table 1 The chemical composition of the Al–5Ti–1B master alloy rods

Batch

Content (weight percentage, wt%) Ti

B

Fe

Si

V

K

1

5.12

1.09

0.10

0.081

0.011

0.028

2

5.09

1.06

0.11

0.090

0.013

0.027

3

5.11

1.05

0.11

0.093

0.021

0.031

Average

5.11

1.07

0.11

0.088

0.015

0.029

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of 0.2 wt% Al–5Ti–1B master alloy rods when the melt temperature was 720 °C, respectively. The third group: four different melt processing temperatures of 690, 720, 750, and 780 °C were selected to add 0.2 wt% Al–5Ti–1B master alloy rods, respectively. After 10 min of stirring and holding, the melt was cast into the mold. After casting, the aluminum ingots were sawed off along the middle axle. The grain structures were observed after grinding, polishing and etching with a strong mixed acid solution (70 ml HCl + 25 ml HNO3 + 5 ml HF) of the half ingots and average grain sizes were measured by quantitative metallography method.

3 Results and Discussion 3.1 Chemical Composition The contents of alloying elements and main impurity elements of the Al–5Ti–1B master alloy rods are listed in Table 1. It indicates that the content of Ti and B elements in the three batches of Al–5Ti–1B master alloy rods are relatively stable, with an average value of 5.11 and 1.07 wt%, respectively. In order to avoid secondary pollution of the aluminum melt which influences the quality of aluminum products, the content of impurity elements Fe, Si, V, and K needs to control. For example, the production of high precision aluminum products such as foil, cans, and PS version the content of impurity elements within Al–5Ti–1B master alloy rods is as lower as better. As shown in Table 1, the impurity elements Fe, Si, V, and K in the Al–5Ti–1B master alloy rods has an average value of 0.11, 0.088, 0.015, and 0.11 wt%, respectively. They are far less than that of the standard Al–5Ti–1B master alloy rods in China. In addition, they are also lower than that of the Al–5Ti–1B master alloy rods produced by the England LSM, Dutch KBW, Spain Aleastur, and South Korea SLM companies.

3.2 Microstructure Figure 1 shows the microstructures and EDX analysis of the Al–5Ti–1B master alloy rods. It consists of TiAl3 and TiB2 particles and without oxides or inclusions. The TiAl3 particles are in block-shaped or plate-shaped and have a uniform average size of 16.7 μm (Fig. 1a, indicated by black arrows). The TiB2 particles are dispersedly distributed in the α-Al matrix with an average size of 0.73 μm, and the maximum size of its agglomerate is less than 5 μm (Fig. 1a and b, indicated by white arrows). Their size, shape, and distribution could take an important effect on the grain refinement performance of the Al–5Ti–1B master alloy rods, and thus the quality of aluminum products. The fine block TiAl3 phase and dispersion of fine granular TiB2 particles could enhance the grain refining effect, while coarse TiAl3 phase and TiB2 particles would reduce it [9]. In addition, the coarse TiB2 particles and its agglomeration

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Fig. 1 Microstructures (a, b) and c, d EDX analysis of the Al–5Ti–1B master alloy rods

are ready to cause defects such as pinhole, tear and scratch on the surface of high precision aluminum products such as foil, cans, and high-grade PS version.

3.3 Grain Refining Performance The as-cast grain structure of the industrial pure aluminum ingot without the addition of the Al–5Ti–1B master alloy rods is shown in Fig. 2. It mainly consists of coarse columnar crystals and a small part of equiaxed grains in the central zone (diameter about 10 mm). The average grain size is large of 2800 μm. In addition, within the surface layer (thickness about 5 mm) relatively smaller columnar crystals are formed due to the rapid cooling of the steel mold. The as-cast grain structures of the industrial pure aluminum ingots after adding different amounts, i.e. 0.05, 0.1, 0.2, and 0.3 wt%, of the Al–5Ti–1B master alloy rods are shown in Figs. 3, and 4 presents the relationship of the average grain size and the additive amount. It indicates the coarse columnar crystals (seen in Fig. 2) are fully transformed into fine equiaxed grains only by a small additive amount of the 0.05 wt% Al–5Ti–1B master alloy rods (as seen in Fig. 3a). Furthermore, the average grain size is significantly decreased with the increase of the additive amount

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Fig. 2 As-cast grain structure of the industrial pure aluminum ingot without the addition of the Al–5Ti–1B master alloy rods

20mm

up to 0.2 wt% (Figs. 3 and 4). However, further increase of the additive amount up to 0.3 wt%, the decrease of the average grain size is limited. By the addition of the 0.2 wt% Al–5Ti–1B master alloy rods, the average grain size of the as-cast industrial pure aluminum ingot has been significantly refined from 2800 μm to 75.7 μm. This demonstrates that the Al–5Ti–1B master alloy rods produced by our research group have great grain refining performance. The effect of holding time after the addition of 0.2 wt% Al–5Ti–1B master alloy rods on the as-cast grain structure of the industrial pure aluminum ingot is shown in Figs. 5, and 6 presents the relationship of the average grain size and the holding time. It indicates that the average grain size is rapidly refined from 2800 μm to 88.3 μm after holding time of 2 min (as seen in Fig. 5a). With the extension of holding time, the average grain size is continually slow decreased and then slowly increases. Even though the holding time prolonging to 120 min, the average grain size is only grown up to 109.5 μm. Therefore, the Al–5Ti–1B master alloy rods possess a fast grain refinement response performance and a high resistance to grain refinement recession. In view of the melt, the temperature could be different when adding the Al–5Ti–1B master alloy rods in various manufacturers of aluminum and its alloy, the response ability to different melt temperatures is also one of the important indexes of grain refining performance [2]. The effect of melt processing temperatures when adding the 0.2 wt% Al–5Ti–1B master alloy rods on the as-cast grain structure of the industrial pure aluminum ingots are shown in Figs. 7, and 8 presents the relationship of the average grain size and the melt processing temperature. It can be seen that fully fine equiaxed grains are obtained at a wide range of melt processing temperature, i.e.,

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(a)

(b)

(c)

(d)

20 mm Fig. 3 The effect of the additive amount of the Al–5Ti–1B master alloy rods on the as-cast grain structure of the industrial pure aluminum ingot: a 0.05 wt%; b 0.1 wt%; c 0.2 wt%; d 0.3 wt% Fig. 4 Relationship between the average grain size of the as-cast industrial pure aluminum ingot and the additive amount of the Al–5Ti–1B master alloy rods

Microstructure and Grain Refining Performance …

(a)

(b)

(c)

(d)

139

20mm Fig. 5 The effect of holding time after the addition of 0.2 wt% Al–5Ti–1B master alloy rods on the as-cast grain structure of the industrial pure aluminum ingot: a 2 min; b 10 min; c 60 min; d 120 min Fig. 6 Relationship between the average grain size of the as-cast industrial pure aluminum ingot and the holding time after the addition of 0.2 wt% Al–5Ti–1B master alloy rods

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(a)

(b)

(c)

(d)

20mm Fig. 7 The effect of melt processing temperature when adding the 0.2 wt% Al–5Ti–1B master alloy rods on the as-cast grain structure of the industrial pure aluminum ingot: a 690 °C; b 720 °C; c 750 °C; d 780 °C

690–780 °C (Fig. 7). The average grain size is almost unchanged at the range of 690–750 °C (Fig. 8). At high melt processing temperature of 780 °C, the average grain size only has a little increase. Thus, it confirms that the Al–5Ti–1B master alloy rods possess a high response ability to melt temperature, i.e., it can obtain a significant grain refining performance at low melt temperature as well as high melt temperature.

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Fig. 8 Relationship between the average grain size of the as-cast industrial pure aluminum ingot and the melt processing temperature when adding the 0.2 wt% Al–5Ti–1B master alloy rods

4 Conclusions (1) The main alloying elements Ti and B are relatively stable with an average value of 5.11 and 1.07 wt%, respectively, and the main impurities Fe, Si, V, and K are relatively low with an average value of 0.11, 0.088, 0.015, and 0.11 wt%, respectively. (2) The TiAl3 particles formed in the Al–5Ti–1B master alloy are in block-shaped or plate-shaped with a uniform average size of 16.7 μm. The TiB2 particles are dispersedly distributed in the α-Al matrix with an average size of 0.73 μm, and the maximum size of its agglomerate is less than 5 μm (3) The optimal addition of the Al–5Ti–1B master alloy rods is 0.2 wt% and could significantly refine the average grain size of the as-cast industrial pure aluminum ingot to 75.7 μm. It possesses a fast grain refinement response, high resistance to grain refinement recession and an adaption for a wide range of melt processing temperature. Acknowledgements We thank the GDAS’ Project of Science and Technology Development (Grants No. 2018GDASCX-0966, 2018GDASCX-0117) and the Science and Technology Planning Project of Guangdong Province, China (Grants No. 2017A070702019, 2017A070701019) for the financial support.

References 1. D. McCartney, Grain refining of aluminium and its alloys using inoculants. Int. Mater. Rev. (1, 34), 247–260 (1989) 2. B.S. Murty, S.A. Kori, M. Chakraborty, Grain refinement of aluminium and its alloys by heterogeneous nucleation and alloying. Int. Mater. Rev. 47(1), 3–29 (2002)

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3. M. Shi-Guang, X. Hui, W. Zhu-Tang, Review and outlook of output of aluminum product and grain refiner requirement in the world. Light Alloy Fab. Technol. 39(10), 1–9 (2011) 4. W. Shun-Cheng, Z. Cai-Jin, Q. Wen-Jun, Z. Kai-Hong, L. Jian-Xiang, Quality evaluation of overseas Al-Ti-B grain refiners. Light Alloy Fab. Technol. 39(6), 11–14 (2011) 5. X. Chen, Q. Ye, J. Li, C. Liu, Y. Yu, Electromagnetic Induction Electric Melting Furnace for Controlling Average Nominal Diameter of TiB2 (TiC) Particle Group in Al-Ti-B (Al-Ti-C) Alloy. CN 201010110166 6. X. Chen, Q. Ye, J. Li, C. Liu, Y. Yu, Purification Method of Al-Ti-B Alloy Melt. CN 201010110046 7. X. Zhang, X. Chen, J. Li, C. Liu, S. Li, Method for Controlling Variable Quantity of Grain Refining Capability of Aluminum-Titanium-Carbon Alloy During Pressure Processing of AluminumTitanium-Carbon Alloy. CN 201010110051 8. YS/T 447.1-2002, Aluminium and aluminium alloys grain refiners: Part 1: Al-Ti-B wire 9. G. Song, D. Shu, W. Lei, H. Yan-Feng, W. Jun, S. Bao-de, Research progress of Al-Ti-B grain refiner. Light Alloy Fab. Technol. 35(12), 7–10 (2007)

Effect of Cu Content on Microstructure and Properties of Al–Mg–Si Alloy Hong-Xiang Li, Shengli Guo, Peng Du and Sheng-Pu Liu

Abstract The effect of Cu content on the microstructure and properties of Al–Mg–Si alloy was studied by optical microscopy scanning electron microscopy, transmission electron microscopy, and tensile test. The state of the heat treatment is directly aged at 165 °C after a solid solution. The results show that with the increase of Cu content, the tensile strength of Al–Mg–Si alloy first increases and then decreases, and the conductivity has no obvious change. When the content of Cu is 0.3%, the tensile strength of Al–Mg–Si alloy is significantly increased to 336 MPa and the electrical conductivity is 49.6% IACS.

1 Introduction Protecting the environment and saving resources is a necessary condition for building sustainable development society. Therefore, the demand for low density, high strength, and high conductivity metal materials is increasing. Aluminum has low density and high conductivity, so aluminum alloy is the preferred material in the industry [1]. It is a common method to add a small amount of Mg and Si into aluminum alloy to produce conductors in the power field [2, 3]. The Al–Mg–Si alloy mainly contains magnesium and silicon, which belong to heat-treated aluminum alloy. The alloy has the advantages of moderate strength, good processing ability, and corrosion resistance, good extrusion performance and so on [4]. Cu element can significantly improve the strength of Al–Mg–Si alloy, but at the same time, it will reduce the corrosion resistance of Al–Mg–Si alloy [5, 6]. On the premise of fixing the content of Mg and Si of Al–Mg–Si alloy, by adding different contents of Cu element, the influence of Cu content on the microstructure, mechanical and electrical properties of Al–Mg–Si alloy was investigated on the basis of the experiment.

H.-X. Li · S. Guo (B) · P. Du · S.-P. Liu General Research Institute for Non-ferrous Metals, Beijing 100088, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_15

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Alloy

Al

Mg

Si

Cu

Mn

A

Balance

0.6

0.75

0.1

0.003

B

Balance

0.6

0.75

0.2

0.003

C

Balance

0.6

0.75

0.3

0.003

D

Balance

0.6

0.75

0.4

0.003

2 Experimental 2.1 Material Al–Mg–Si alloy was melted with graphite crucible in a 7.5 kW electric furnace. When the temperature raised to 760 °C, the Al–Cu master alloy was added into the water-cooled copper mold and cast into a round ingot with a diameter of 120 mm. Al–Mg–Si alloy ingots with Cu mass fraction of 0.1, 0.2, 0.3, and 0.4%, respectively, were named A, B, C, and D alloys. The contents of the main chemical components Mg and Si remained unchanged. The specific components are shown in Table 1. The aluminum alloy ingot was extruded on a 1250 t extruder to form a rectangular aluminum strip with a length of 60 mm and a width of 6 mm. Extrusion ratio is 32; The solid solution process is (540 °C, 2 h), and the aging process is (165 °C, 0–24 h).

2.2 Characterization Metallographic samples were smoothed and polished with water sandpaper for chemical etching (the corrosion solution: 17mLHF + 3mLHCl + 5mLHNO3 + 75mLH2 O). The microstructure of the alloy was observed by an Olympus-GX51 optical microscope (OM). The hardness test was carried out on the Victorinox hardness tester. Each sample was measured with 7 points and the average value was obtained. The microstructure of the sample was observed in the JEM-3010 transmission electron microscope. The fracture morphology was observed by PHILIPS-XL30 scanning electron microscope, and the tensile properties were measured by the WEW-1000B universal material test machine, and the strain rate was 1 mm/min. The electrical conductivity of the sample was measured by FD-101 eddy current conductivity meter.

Effect of Cu Content on Microstructure and Properties …

Hardness/HV

Fig. 1 The age-hardening curves of Al–Mg–Si alloy with different Cu contents

145

130 125 120 115 110 105 100 95 90 85 80 75 70 65 60 55

A Alloy B Alloy C Alloy D Alloy

0

2

4

6

8

10

12

14

16

18

20

22

24

26

28

Ageing time/h

3 Results and Discussion 3.1 Age-Hardening Curve It can be seen that at the same aging time, the hardness of the alloy first increased and then decreased with the increase of Cu content (Fig. 1). The 0.3 mass% Cucontaining alloy shows the highest peak hardness, reaching to 124.6 HV, and the peak hardness time is reduced by about 2 h. This phenomenon is different from the previous results reported in the previous literature. The increase of the Cu content delays to achieve the peak hardness by single stage and two stage [7]. The study of 6022 aluminum alloy shows that the peak hardness of the alloy with different Cu content appears the same time [8]. That is, the difference of Cu content will not affect the time of the peak. In the absence of aging, the hardness of C alloy is close to that of D alloy, which is higher than that of A and B alloy. With the prolongation of aging time, the hardness of C alloy is always at the highest value of the hardness of the four alloys. The hardness of D alloy is slightly lower than that of C alloy. This may be due to the uneven diffusion of CuAl2 phase produced by aging in the alloy when Cu content reaches 0.4%. The above results show that the Cu content is not more than 0.4%, which is beneficial to increase the peak hardness of the alloy, reduce the peak hardness time, and increase the aging hardness of the alloy as a whole.

3.2 Tensile Strength and Electrical Conductivity With the increase of Cu content, the tensile strength increases first and then decreases, but the electrical conductivity does not fluctuate significantly (Fig. 2). The tensile

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Fig. 2 Effect of different Cu contents on the properties of Al–Mg–Si alloys aged at 165 °C for 6 h

350 340 330

Strength/MPa

320 310 300 290 280 270 260 250 0.1

0.2

0.3

0.4

0.3

0.4

ω (Cu)% 55 54 53

Conductivity/% IACS

52 51 50 49 48 47 46 45 44 43 42 41 40 0.1

0.2

ω (Cu)%

strength reaches to 336 MPa at 0.3% Cu-containing alloy, and the conductivity is 49.6% IACS. This is mainly attributed to Cu played a role in solid solution strengthening, which will increase the tensile strength [9]. Comparing the tensile strength and electrical conductivity of four kinds of aluminum alloys with different Cu contents, it is found that under the same processing, the proper addition of Cu content is beneficial to increase the tensile strength of the sample, but has little effect on the conductivity. The radius of the Cu atom is close to that of Al atom, causing the small lattice distortion of the aluminum. According to the theory of metal conduction, when the lattice distortion is very small, it has little effect on the conductivity.

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3.3 Optical Microstructure The as-cast microstructure of alloy with 0.1% Cu content is mainly coarse dendrites with large dendrite spacing and solute atoms are evenly distributed in the grains (Fig. 3a). As can be seen from Fig. 3b and c, with the increase of Cu content, the spindle of the dendrite becomes finer, and the spacing is narrowed, which is composed of fine dendrites and the solute atoms are enriched in the grain boundary and near the boundary. This is mainly due to the formation of CuAl2 [6] in the aluminum alloy after the Cu addition, which is distributed near the grain boundaries and inhibits grain boundary migration and hinders grain growth. The dendrite size of D Alloy is uneven, which is composed of coarse dendrites and fine dendrites. This may be due to the uneven dispersion of the CuAl2 phase formed in the alloy when the Cu content increases to 0.4%. The grain is broken along the direction of deformation and the grain boundaries become clear and straight (Fig. 3). The alloy structure changes from the dendritic grain to a uniform equiaxed grain, eliminating dendrite segregation, and improving the mechanical and electrical properties of the alloy [10]. The average grain size of A, B, and C alloys is about 130, 60, and 30 μm, respectively (Fig. 3). It is obvious that the grain size decreases with the increase of Cu content. The grain size of the alloy decreases gradually, and the grain boundaries become more numerous. The resistance of free electrons passing through the grain boundaries increases, and the conductivity decreases slightly [11]. The grain in Fig. 3h is composed of large size grain and small size grain, which may be caused by the dynamic recrystallization in the process of extrusion, because the dispersion of CuAl2 phase is not uniform. The extruded aluminum alloy ingot can make the structure more compact, reduce the porosity of ingot, and refine the microstructure [12, 13].

3.4 Tensile Fracture Figure 4 shows the tensile fracture morphology of Al–Mg–Si alloy with different Cu contents after solid solution (540 °C, 2 h) and aging (165 °C, 6 h). The aluminum matrix is a face-centered cubic structure. The tensile fractures are generally divided into three types, namely, slip band cracking, intergranular cracking, and dimple-type cracking. With the increase of Cu content, the plasticity of the alloy fracture becomes more and more obvious, especially when the Cu content is 0.3%, the dimple formed by the tensile fracture becomes deeper and more (Fig. 4c). It is indicated that the plasticity of the material increases after the increase of Cu content, which is consistent with the increase in elongation as increasing Cu content. Figure 4d shows that the fracture surface has a dimple, but the dimple is relatively shallow and small, which is consistent with the result of the strength reduction of D alloy.

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Fig. 3 The microstructures of as-cast a A alloy, b B alloy, c C alloy, and d D alloy, and extruded and solutionized e A alloy, f B alloy, g C alloy, and h D alloy

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Fig. 4 Fracture morphology of alloy with different Cu contents: a A alloy; b B alloy; c C alloy; d D alloy

3.5 HRTEM Observation Figure 5 shows the HRTEM image of A and C alloy after solid solution (540 °C, 2 h) aging (165 °C, 6 h). As can be seen from Fig. 5a, some fine spherical GP regions are dispersed in the matrix of A alloy, and a small amount of pre-β , only a small number of short rod-like precipitates. There are two kinds of GP zone strengthening in the C alloy. Besides the GP region formed by Mg and Si elements, Al and Cu elements also form the GP region, which makes the strength increase. Compared with the A alloy, the formation of the GP region of the Al and Cu elements is a common lattice (Fig. 5c). The lattice distortion is smaller and the additional scattering of the electronic motion is weakened, which leads to a little decrease in electrical conductivity [14].

4 Conclusions With the increase of Cu content, the peak hardness of Al–Mg–Si alloy can be increased and the time to peak hardness can be reduced. Cu addition into the Al–Mg–Si alloy can form the GP region, which increased strength remarkably while

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Fig. 5 HRTEM of alloys with different Cu contents: a A alloy; b C alloy; c the red area in Fig. 5b

the electrical conductivity changed a little. This is because the GP region formed by Al and Cu had a little scattering effect on electrons. When the content of Cu was 0.3%, the tensile strength of Al–Mg–Si alloy was significantly increased to 336 MPa, and the conductivity was 49.6% IACS.

References 1. X. Sauvage, E.V. Bobruk, Y. Nasedkina et al., Optimization of electrical conductivity and strength combination by structure design at the nanoscale in Al–Mg–Si alloys. Acta Mater. 98, 355–366 (2015) 2. F. Kiessling P. Nefzger, F. Nolasco, Overhead Power Lines: Planning, Design, Construction (Berlin: Springer, 2003) 3. G.E. Totten, D.S. MacKenzie (eds.), Handbook of Aluminium. in Alloy production and Materials Manufacturing, vol. 2 (New York: Marcel Dekker, 2003)

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4. J. Hirsch, Aluminium Alloys for Automotive Application: Materials Science Forum (1997) 5. J.X. Zhang, A.H. Gao, Influence of trace of Cu on microstructure and properties of 6063 aluminum alloy. Mater. Heat Treat. 38(22), 45–47 (2009) 6. H.-X. Gao, X.-Y. Zhang, X. Huang et al., The effect of Zr, Cu elements and rare earth Ce on microstructure and property of the electrical round rod. J. Funct. Mater. 46(3), 3073–3076 (2015) 7. Y.N. Liu, J.H. Chen, M.J. Yin et al., The influences of natural ageing and Cu addition on the age hardeni ng behavior of AlMgSi(Cu) alloys. J. Chinese Electron Microsc. Soc. 29(3), 280–286 (2010) 8. W.F. Miao, D.E. Laughlin, Effects of Cu content and preaging on precipitation characteristics in aluminum alloy 6022. Metall. Mater. Trans. A 31(2), 361–371 (2000) 9. X.J. Shang, Q.B. Liu, P. Xu et al., Effects of copper and rare Earth elements on properties of aluminum electrical round bars. Nonferrous Met. Eng. 8(1), 16–19 (2018) 10. X.Y. Zhang, H. Zhang, X.X. Kong et al., Microstructure and properties of Al-0.70Fe-0.24Cu alloy conductor prepared by horizontal continuous casting and subsequent continuous extrusion forming. Trans. Nonferrous Met. Soc. China 25(6), 1763–1769 (2015) 11. M. Suzuki, T. Kimura, J. Koike et al., Strengthening effect of Zn in heat resistant Mg–Y–Zn solid solution alloys. Scripta Mater. 48(8), 997–1002 (2003) 12. F. Li, X.D. Liu, W.Y. Wang et al., Effect of squeeze casting on microstructure of A356 alloy. Foundry 57(4), 347–349 (2008) 13. X.M. Zhang, L.H. Hao, D.M. Jiang et al., An investigation on tensile fracture of Al-Mg-Si alloys. Mater. Eng. (5), 35–36 (1996) 14. H.F. Liu, N.Z. Sun, X.L. Dai et al., The measurement and researchment of temperature of GP domain in the procedure of age-hardening. Hot Work. Technol. 2, 81–85 (1988)

Corrosion Behavior of Inconel 625 Alloy in Na2 SO4 –K2 SO4 at High Temperature Yuan-Jun Ma, Yutian Ding, Jian-Jun Liu, Yu-Bi Gao and Dong Zhang

Abstract The corrosion behavior of Inconel 625 alloy at 800 and 900 °C in the molten salt of 75 wt% Na2 SO4 –25 wt% K2 SO4 was studied. The corrosion mechanism of Inconel 625 alloy after hot corrosion is mainly alkaline melting mechanism. The Cr2 O3 on the surface of the alloy dissolved in molten salt as Na2 CrO4 , leading to the loss of the protective oxide layer on the alloy surface. With the decomposition of Cr2 O3 on the surface of the alloy, Cr-depleted region appeared at the interface of the alloy matrix/corrosion layer, which inhibited the growth of the Cr2 O3 oxide layer and resulted in the discontinuous oxide layer. This caused O and S to invade the substrate and corrode the alloy matrix. When the alloy was corroded at 800 °C for 120 h, the corrosion rate was about 3 mg/cm2 . The corrosion layer was relatively complete, and it mainly consisted of flaky Cr2 O3 and spinel-like NiCr2 O4 . When the alloy was corroded at 900 °C for 120 h, the corrosion rate was about 6 mg/cm2 , and obvious shedding and faults appeared in the corrosion layer, which were mainly divided into three layers: the outer layer was composed of NiCr2 O4 and NiO; the middle layer was a dense Cr2 O3 ; the inner layer was composed of sulfides (Cr2 S3 and Ni3 S2 ), oxides (Cr2 O3 and NiO), telluride, etc. Through thermodynamic calculation and analysis, it was found that SO2 decomposed from molten salt under high temperature could cause severe corrosion to the alloy. The main corrosion products were Cr2 O3 , Cr2 S3 , NiO, and Ni3 S2 .

1 Introduction Ni-based superalloy Inconel 625 (IN625) has been one of the most widely used in aerospace, petrochemical, chemical and marine applications due to its excellent Y.-J. Ma · Y. Ding (B) · J.-J. Liu · Y.-B. Gao State Key Laboratory of Advanced Processing and Recycling of Non-ferrous Metals, Lanzhou University of Technology, Lanzhou 730050, China e-mail: [email protected] D. Zhang State Key Laboratory of Nickel and Cobalt Resources Comprehensive Utilization, Jinchuan Group Co., Ltd., Jinchang 737100, China © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_16

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corrosion and oxidation resistance, as well as its high yield strength, creep strength, and fatigue strength [1, 2]. The service environment of the alloy is very complex, such as sulfur, carbon, and chlorine. Moreover, the deposits of metallic sulfates and chlorides, such as Na2 SO4 , may also accumulate on the exposed surface. Under certain conditions, these deposited salts give rise to severe hot corrosion attack and accelerate its degradation [3–5]. Therefore, many researchers have carried out a great deal of research on its high-temperature corrosion resistance, and proposed the general mechanism of high-temperature corrosion [4–8]. However, due to the influence factors of hot corrosion is very complicated, such as alloy composition, preparation process, the composition of sediments, the sediment amount and corrosion gas composition, flow rate, temperature and corrosion mode are important factors affecting corrosion [9–12]. Moreover, the corrosion resistance of the alloy under different corrosion conditions will also vary greatly. Therefore, it is an important job to evaluate the corrosion resistance of the alloy by using the simulated environment of the alloy. In coal combustion equipment, Na2 SO4 and K2 SO4 were commonly deposited. However, the hot corrosion behavior of alloy in molten sulfate at high temperature (800 and 900 °C) was less reported, especially for Inconel 625 alloy. Therefore, this paper mainly studied the hot corrosion behavior of Inconel 625 alloy in 75 wt%Na2 SO4 –25 wt%K2 SO4 mixed molten salt at 800 and 900 °C. Through the analysis of the composition, morphology, and distribution of corrosion products, so as to provide a theoretical reference for the application of Inconel 625 alloy in high-temperature corrosion environments.

2 Experimental The material used in the test is cast Inconel 625 alloy. The alloy is produced by vacuum induction melting and electroslag remelting (VIM + ESR). The main chemical composition is shown in Table 1. The alloy was cut into a round specimen with dimensions of 14 mm × 5 mm by the linear cutting method. The specimens were polished with abrasive paper, and then 5 min were cleaned, respectively, in acetone and alcohol, and dried. The salt mixture was made from anhydrous Na2 SO4 and K2 SO4 . The Na2 SO4 and K2 SO4 purity were greater than 99.0 wt%. Then, 15.0000 g Na2 SO4 and 5.0000 g K2 SO4 were weighed by an electronic analytical balance (FA1004) with an accuracy of 0.1 mg, mixed extremely well and put in one corundum crucible. Add a small amount of distilled water to the mixture, then weigh the prepared sample with an electronic balance and place it in the mixed salt. Put the crucible with samples and mixed salt into the drying

Table 1 Inconel 625 alloy chemical composition (mass fraction, %) Ni

Cr

Nb

Co

Mo

C

Al

Ti

Fe

Si

S

Bal.

21.77

3.75

0.19

8.79

0.042

0.21

0.40

3.68

0.12

0.0006

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(a)

(b)

3

4

2 1

corundum 5

crucible

Fig. 1 a Experimental setup and b Experimental equipment. 1: sample: 2: molten salt; 3: thermocouple; 4: silicon carbide rod; 5: temperature controller

oven (101-1A), drying at 100 °C for 30 min, as shown in Fig. 1a. During these time, the box type resistance furnace (SX-G18123) has been heated to 800 and 900 °C at a heating rate of 5 °C/min. Then put the crucible with the mixture of salt and sample into the resistance furnace directly, as shown in Fig. 1b. After 10, 20, 40, 60, 80, 100, and 120, the crucibles with the salt samples were removed from the furnace and cooled to room temperature. There are three parallel patterns under each corrosion condition. Then boil the sample in distilled water for 30 min to remove the mixture of salt which is attached to the sample, then wash it with alcohol and weigh it after drying. The following are analysis methods of surface corrosion: The phase of the sample was examined by X-ray diffractometer (D8ADVANCE). The morphology and composition of samples were analyzed by field emission scanning electron microscope (SEM) with Oxford INCA spectrometer and its attachment EDAX.

3 Results and Discussion 3.1 Hot Corrosion Behavior Figure 2 shows the change of mass loss of Inconel 625 alloy after corroded in a mixed salt of Na2 SO4 and K2 SO4 at a molar ratio of 3:1 at 800 and 900 °C. When the corrosion temperature is 800 °C, it belongs to low-temperature hot corrosion. The corrosion temperature does not reach the melting point of the mixed salt, and the mixed salt is in a solid state. But during the decomposition process, the eutectic phase Na2 SO4 –NiSO4 (melting point 671 °C) and Na2 SO4 –CoSO4 (melting point

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Fig. 2 Mass change of Inconel 625 alloy after corrosion

mass change / (mg/cm2)

8

6

4

2

0

800 900

0

20

40

60

time / h

80

100

120

565 °C) with a low melting point is formed. This leads to the occurrence of a small amount of liquid phase, which causes the accelerated corrosion of the material [13]. Since the amount of molten salt is very small, the corrosion process is also very slow, and the corrosion rate is about 3 mg/cm2 . When the corrosion temperature is 900 °C, the corrosion is high-temperature hot corrosion. The corrosion temperature is higher than the melting point of the mixed salt (835 °C), the mixed salt is in a molten state. Therefore, the corrosion rate of the alloy is faster, and the corrosion rate of the alloy at 900 °C is 6 mg/cm2 . Figure 3 shows the XRD patterns of the as-cast Inconel 625 alloy after hot corrosion under different conditions. The X-ray diffractometer uses CuKα as the diffraction source. The operating tube voltage and tube current are 40 kV and 150 mA, respectively. The scanning range is 10°–90°. Figure 3a shows the XRD pattern of the as-cast Inconel 625 alloy after 120 h oxidation at high temperature. It can be seen that the pure oxidation products of Inconel 625 alloy are mainly Cr2 O3 , NiCr2 O4 , NiO and MoO. Figure 3b shows the XRD pattern of mixed salts around the sample after hot etching in mixed salt at 800 °C. The mixed salts are Na2 SO4 and K2 SO4. The yellow mixed salt is Na2 CrO4 (shown in the small panel in Fig. 3b). Because of the low-temperature hot corrosion and the amount of molten salt produced, the surface of the alloy is surrounded by a dense solid mixed salt layer. Oxygen is difficult to diffuse into, so excess Na2 O will react with Cr2 O3 to form Na2 CrO4 , dissolved in molten salt. Figure 3c is the analysis of the phase structure of Inconel 625 corrosion film by GIXRD. The grazing angle is 2.5° and the scanning range is 10°–90°. It can be seen that the composition of the corrosion products under different etching times does not change much. The oxides are mainly such as Cr2 O3 , NiCr2 O4 , NiO, and MoO, as well as sulfides such as Cr2 S3 and Ni3 S2 . With the increase of corrosion time, the relative intensity of the diffraction peak of Cr2 O3 decreases, there is an

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Fig. 3 XRD spectrum of Inconel 625 alloy after hot corrosion. a oxidation; b mixed salt; c corrosion at 900 °C at different times; d corrosion at different temperatures for 120 h

obvious diffraction peak of NiCr2 O4 , and the relative intensity of the diffraction peak of NiCr2 O4 increases. The duration of incubation is also increasing. Figure 3d is an XRD pattern of a hot-etched sample of as-cast Inconel 625 alloy in mixed salt at different temperatures. It can be found that at 800, and 900 °C, the corrosion products are basically the same, mainly Cr2 O3 , NiO, and NiCr2 O4 .

3.2 Analysis of Corrosion Product Analysis of Corrosion Layer Figure 4 is a SEM diagram of Inconel 625 alloy after hot corrosion of 800 °C at different times. As can be seen from Fig. 4a and b, the alloy shows different features after the hot corrosion for 20 and 120 h. They all form a layer of uneven dense corrosion layers consisting of particles of different sizes, the size of the particles on the corrosion surface of the sample which is obviously corroded for 120 h is larger. Figure 5 is a mass fraction diagram of the elements at different points in the SEM diagram of Inconel 625 alloy after hot corrosion at 800 °C. As can be seen from Fig. 5, the outer layer of corrosion layer is mainly composed of Ni, Cr, and O. Combined with XRD analysis, Cr2 O3 , NiO, and NiCr2 O4 can be identified.

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Fig. 4 SEM diagram of Inconel 625 alloy after hot corrosion at 800 °C. a 20 h; b 120 h; c 60 h; d 60 h; e 60 h Fig. 5 The mass fraction of the elements in the SEM diagram for Inconel 625 alloy after hot corrosion at 800 °C

Nb Fe Si

Mass fraction/(%)

Ni Cr S O

Point

Perform EDS energy spectrum analysis of points 1, 2, 3, and 4 in Fig. 4a and b. According to XRD analysis, the surface of Inconel 625 alloy etched at 800 °C for 20 and 120 h is mainly composed of larger spinel NiCr2 O4 and smaller spinel Cr2 O3 . However, NiCr2 O4 is easily detached from the surface of the sample because its stability is not as good as that of Cr2 O3 . It is evident from Fig. 4c that partial oxidation layers are preferentially damaged (as shown in arrow 5, 6). EDS spectroscopy analysis (as shown in points 5, 6 in the Fig. 5) found that the regions preferentially damaged in the oxide layer were rich in Nb elements. This heterogeneous corrosion is due to the segregation of Nb or Nb compounds in nickel-based alloys, and the stability of these Nb compounds is not preferentially high in the matrix, so they are preferentially eroded during corrosion.

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As corrosion proceeds, a clear crevice occurs where some of the oxide layers are preferentially corroded and destroyed (as shown in the arrow in Fig. 4d). EDS analysis (Fig. 5, point 8) revealed that the S element appeared here, while the mass fraction of other elements were essentially unchanged compared to the outer oxide layer (as shown in point 7, 8 in the Fig. 5). This indicates that the compounds of S and S enter the inner layer of the alloy matrix through the damaged oxide layer and leave corrosion pits. On the other hand, it is also indicated that the dense Cr2 O3 oxide is the main cause of the sulfur corrosion resistance of the alloy. During the corrosion process, the Cr element diffused to the surface of the matrix and the Cr2 O3 gradually decomposed, thus consuming the Cr element in the matrix, resulting in the decrease of its content. The sulfide in the corrosion products can reduce the viscosity of the oxide film to the alloy and promote the falling off of the oxide film [14]. The surface oxide layer will become loose or flaking (as shown in the arrow in Fig. 3e), and corrosion holes appear in the inner layer (as shown in Fig. 3e). When the content of Cr in the matrix is relatively small and the content of Ni is relatively large, because of the little affinity between Ni with O and Cr, the complex oxide NiCr2 O4 will form. At the same time, some Ni elements will form NiO on the metal surface, which is consistent with the XRD results of Fig. 3c. Generally speaking, low-temperature hot corrosion leads to the formation of eutectic with a low melting point, and the corrosion mechanism is the same as that of high-temperature hot corrosion. However, due to the limited amount of molten salt, the corrosion degree of low-temperature hot corrosion is not high, and the damage of oxide layer on the surface of the alloy is limited. Figure 6 is the SEM diagrams of Inconel 625 alloys after high-temperature corrosion at 900 °C for different times. As can be seen from the diagram, the surface of the alloy is uneven after the hot corrosion of 120 h (as shown in Fig. 6a), and a clear exfoliation of the oxide layer occurs (as shown in area E in the Fig. 6). By looking at different levels of spalling, it can be seen that the corrosion layer consists of many layers. Figure 7 is a mass fraction diagram of the elements at different points in the SEM diagram of Inconel 625 alloy after hot corrosion. As can be seen from Fig. 7, the corrosion layer consists of several layers, and the content of the elements at different points has changed greatly. It can be seen from Fig. 6a that the “A” zone is the complete oxidation zone, the “C” zone can find the obvious oxide layer corrosion after spalling, and the “B” area is the obvious corrosion hole. After SEM and EDS analysis, the results of corrosion of Inconel 625 alloy at 800 and 900 °C are similar. The loose oxide of the surface layer gradually exfoliates and decomposes, while the inner layer forms a dense oxide layer again. In Nb segregation, it is impossible to form an effective oxide protective film in time. S reacts easily with the alloy matrix, resulting in severe corrosion. As the corrosion proceeds, the oxide is broken down as it is corroded at 900 °C. The outer surface is changed from loose flaky oxide to compact oxidized layer composed of blocks. In addition, the oxide in the outer and middle layers is loose, and the protection for the matrix is poor. The bottom oxide is dense and has good protection for the matrix. On the whole, there is a tendency of decomposition and diffusion to outward. By comparison, it is found that compounds with Ni and Nb elements

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Fig. 7 The mass fraction of the elements in the SEM diagram of Inconel 625 alloy after hot corrosion at 900 °C

Mass fraction/ (%)

Fig. 6 SEM diagram of Inconel 625 alloy after hot corrosion at 900 °C. a 120 h; b 20 h; c 80 h; d 120 h

Point

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Fig. 8 The cross-section SEM diagram of Inconel 625 alloy after hot corrosion. a 800 °C; b 900 °C Table 2 The element content of section SEM of Inconel 625 alloy after hot corrosion

Element

Ni

Cr

O

Nb

Content A (wt%)

46.04

22.62

14.24

Content B (wt%)

53.47

20.41

5.66

1.23

Content C (wt%)

64.03

19.88

2.1

2.91

Mo

Fe

1.84 8.99

2.09

are mainly formed in places, where Nb elements are abundant, around the pits are mainly oxides of Ni and Cr (as shown in point 7, 10, 12 in Fig. 7). In combination with Figs. 4, 6, and 7, it is also found that the presence of Si elements in the corrosion layer of the alloy matrix is beneficial to the corrosion resistance. According to the existing research results, [15] shows that the composite oxides of Si usually exist in the boundary between the matrix of the near alloy and the oxide film. It is combined with chromium oxide to increase the density of the oxide film and improve the oxidation resistance and thermal corrosion resistance of the alloy. Section Morphology and Composition Analysis of Corrosion Layer Figure 8 is the sectional morphology of the Inconel 625 alloy after hot corrosion at 800 and 900 °C, respectively. It can be seen from Fig. 8a that the alloy has an obvious oxidation layer on the surface of the alloy after hot corrosion at 800 °C, and there is no obvious preferential corrosion. It can be found that the corrosion layer is divided into three layers, the upper layer is about 110 μm thick oxide layer (as shown in A region in the Fig. 8a), the middle layer is about 15 μm thick corrosion layer (as shown in B region in the Fig. 8a), the lower is the alloy matrix (as shown in C area in the Fig. 9a). Combined with Table 2, it is found that the corrosion layer is mainly composed of NiCr2 O4 , NiO, and Cr2 O3 . Figure 9 is the sectional morphology of Inconel 625 alloy after hot corrosion at 800 and 900 °C, respectively. From Fig. 9, it is found that a layer of Cr2 O3 oxide film will be formed on the outer surface of the alloy after corrosion, but the integrity of the

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Fig. 9 The cross-section SEM diagram of Inconel 625 alloy after hot corrosion. a 800 °C; b 900 °C

Cr2 O3 oxide film is obviously destroyed as the corrosion temperature increases from 800 to 900 °C. After corrosion at 900 °C, the aggregation of Cr elements on the surface of the alloy can be found from the distribution energy spectrum of the Cr element, but the distribution is uneven. The distribution of O elements shows that the oxide film on the surface of the alloy is discontinuous and is seriously damaged. The poor Cr zone appears below the surface, indicating that at high temperature, the Cr element near the corrosion layer is consumed, and the internal Cr element cannot spread to the surface in time, which leads to the failure to form a continuous and dense Cr2 O3 oxide film. When the surface Cr2 O3 oxide film is destroyed at high temperature, the corrosion element S will diffuse into the alloy matrix through the incomplete oxide layer and corrode the alloy matrix, as shown in the distribution diagram of the S element. When the oxide layer is destroyed, the S element is aggregated below it. In combination with Figs. 4, 5, 6, and 7, it is determined that a dense Cr2 O3 film is formed on the outer surface of the alloy sample. In places where Cr2 O3 is not covered, molten salts react with the substrate to produce Cr3 S4 , and oxides of Ni, Cr, and Nb. Because the alloy is selective oxidation, the Cr element is diffused in the process of oxidation at high temperature and forms a compact oxide film with O, so as to achieve the effect of corrosion resistance. The diffusion of Cr results in a decrease in the content of Cr in the matrix. Therefore, once the corrosion material breaks through the reaction between the oxide layer and the substrate, it is difficult to form a dense oxide film, so as not to have better corrosion resistance.

3.3 Mechanism Analysis of Hot Corrosion In the hot corrosion of nickel base alloy, the corrosion resistance of the alloy is due to the rapid formation of Cr2 O3 in the alloy, the oxide film of Cr hinders the reaction and reduces the activity of oxygen ion in molten salt [16], so the corrosion

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product contains Cr2 O3 . In this paper, the corrosion of high-temperature molten salt in Inconel 625 alloy is mainly divided into two parts. The first is the formation and decomposition of the oxide film on the surface of the alloy, and the second is the corrosion of the matrix. The corrosion resistance of Inconel 625 alloy is also mainly due to the hindering reaction of the Cr2 O3 oxide film. With the increase of corrosion time, the content of Cr in the alloy is limited. It would lead to poor Cr area in the matrix, the salt on the sample surface reacts with the substrate in the alloy which is not covered by the Cr2 O3 oxide layer and forms other discontinuous oxides, such as Ni, NiO, and Nb oxides. These reactions can cause the O and S to penetrate into the alloy matrix and react with the alloy matrix. Moreover, NiO will react with Cr2 O3 to form spinel material NiCr2 O4 , which can reduce the protection of the Cr2 O3 oxide layer to the alloy. In addition, the Cr of the alloy surface will expand when the Cr2 O3 film is formed. However, the expansion coefficient of the formed oxide film is not consistent with that of the alloy matrix. Therefore, with the increase of corrosion time, the oxide film will crack and fall because of the stress increasing with the alloy matrix, which also makes the corrosion and oxidation elements in the salt film diffuse to the alloy matrix [16]. All these factors aggravate the corrosion of Inconel 625 alloy. In addition to these factors, the most important is the dissolution of the oxide film, and the most important is the acid-alkali fusion model theory. The reaction equation is 9 3 1 1 Na2 SO4 + Ni + Cr2 O3 + O2 → Na2 CrO4 + 3NiO + Ni3 S2 2 2 4 2

(1)

The main factor that is known to resist oxidation and hot corrosion is the Cr element in the material. However, the corrosion reaction produces a low melting point Ni·NiS eutectic (645 °C), sulfide, or complexing salt. These corrosion products prevent the alloy from forming a protective oxide film, which is the most important cause of the increased corrosion. Because of the nonmetallic nature of the S and O elements and the difference in metallicity between Ni and Cr, the analysis of the resulting substances should have priority from a thermodynamic point of view. Figure 10 is a standard Gibbs free energy and temperature diagram for some metals, oxides, and sulfides. The lower the standard Gibbs free energy is, the easier it is to generate, so the more stable of these corrosion products is Cr2 O3 , NiO. This is also the principle that Cr elements protect against Ni-based alloys, which are preferentially oxidized or vulcanized by Ni elements to protect the Ni matrix. It is also shown that the priority of oxidation and sulfidation is generally the easiest to oxidize, followed by vulcanization. Figure 11 is a hot corrosion mechanism diagram of Inconel 625 alloy at high temperature. Hot corrosion is the formation of oxide layer, and then an alkaline environment, the oxide layer is alkaline dissolved and destroyed, resulting in matrix corrosion by S. In this process, the Cr element in the γ matrix diffuses to the surface of the substrate to form a new Cr2 O3 oxide layer, which is then dissolved in an alkaline

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Na2 CrO4 MoO3 Na 2O, NiO, Ni 3S2 NiCr2O4

CrS Ni 3 S2

Na 2MoO4 SO2 NiO, Ni3S2

Cr2O3 , Cr2S3

Na 2CrO4 NiO, Ni 3S2

Cr2 O3, SO2

T/ C

T/ C

Fig. 10 The temperature and free energy diagram of hot corrosion reaction of Inconel 625 alloy

Fig. 11 The mechanism of high-temperature hot corrosion of Inconel 625 alloy

environment. With the reaction, there will be the occurrence of internal sulfidation and oxidation.

4 Conclusion (1) During the hot corrosion, the Cr2 O3 on the corrosion layer of the alloy surface has been dissolved in molten salt in the form of Na2 CrO4 . The corrosion rate of the alloy after corrosion at 800 and 900 °C for 120 h is 3 and 6 mg/cm2 , respectively. (2) Under the low-temperature corrosion, Inconel 625 alloy has only a small amount of molten salt, and the corrosion layer is more complete. They are mainly flaking Cr2 O3 and spinel-like NiCr2 O4 . Under the high-temperature hot corrosion, the corrosion layer appears obviously falling off and faults. They are mainly divided into three layers: the outermost layer is NiCr2 O4 and NiO; the second layers are Cr2 O3 ; the third layers are Cr2 O3 , NiO, Ni3 S2, and Cr2 S3 . (3) The main reason for the decomposition and spalling of Cr2 O3 in the corrosion process of Inconel 625 alloy is that the lack of Cr on the surface of Inconel 625 alloy hinders the growth of Cr2 O3 oxide layer, which results in the formation of the discontinuous oxide layer. Moreover, the presence of molten salt on the alloy

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causes chemical dissolution of Cr2 O3 , which is mainly an alkaline dissolution model. (4) With the increase of corrosion temperature from 800 to 900 °C, the corrosion layer of Inconel 625 alloy is more obvious and the oxide film damage is more serious. As the corrosion time increases from 10 to 120 h, the outer layer of the corrosion layer gradually becomes loose and the sheet Cr2 O3 increases. Generally speaking, the corrosion temperature (800 and 900 °C) has a greater effect on the corrosion resistance of alloys in the short time (10–120 h). (5) In the process of corrosion, O and S are easy to invade the alloy matrix due to the existence of poor Cr zone in the alloy, which leads to internal oxidation and internal vulcanization. Acknowledgements The financial support by the National Natural Science Fund (51661019), Gansu Province Major Science and Technology Special Project (145RTSA004), National Key Laboratory of Nickel and Cobalt Resource Comprehensive Utilization (301170503).

References 1. G.P. Dinda, A.K. Dasgupta, J. Mazumder, Laser aided direct metal deposition of Inconel 625 superalloy: microstructural evolution and thermal stability. Mater. Sci. Eng. A 509, 98–104 (2009) 2. E. Mohammadi Zahrani, A.M. Alfantazi, High temperature corrosion and electrochemical behavior of Inconel 625 weld overlay in PbSO4 -Pb3 O4 -PbCl2 -CdO-ZnO molten salt medium. Corros. Sci. 85, 60–76 (2014) 3. S.J. Zinkle, G.S. Was, Materials challenges in nuclear energy. Acta Mater. 61(3), 735–758 (2013) 4. P.S. Sidky, M.G. Hocking, The hot corrosion of Ni-based ternary alloys and superalloys for application in gas turbines employing residual fuels. Corros. Sci. 27(5), 499–530 (1987) 5. D. Kim, H.J. Lee, C. Jang, D.J. Yoon, Corrosion characteristics of Ni-base superalloys in high temperature steam with and without hydrogen. J. Nucl. Mater. 441(1–3), 612–622 (2013) 6. N. Eliaz, G. Shemesh, R.M. Latanision, Hot corrosion in gas turbine components. Eng. Fail. Anal. 9(1), 31–43 (2002) 7. K. Misraa, Studies on the hot corrosion of a nickel-base superalloy, Udimet 700. Oxid. Met. 25(3–4), 129–161 (1986) 8. W. Kai, C.H. Lee, T.W. Lee, C.H. Wu, Sulfidation behavior of Inconel 738 superalloy at 500–900 °C. Oxid. Met. 56(1/2), 51 (2001) 9. R.D.K. Misra, R. Sivakumar, Effect of NaCl vapor on the oxidation of Ni-Cr alloy. Oxid. Met. 25(1/2), 83 (1986) 10. S.R. Kameswa, The role of NaCl in the hot-corrosion behavior of nimonic alloy 90. Oxid. Met. (1/2), 33 (1986) 11. L. Jintao, Y. Zhen, L. Yan et al., Effect of alloying chemistry on fireside corrosion behavior of Ni–Fe-based superalloy for ultra-supercritical boiler applications. Oxid. Met. 12, 1–13 (2017) 12. H. Cui, J.S. Zhang, Y. Murata et al., Hot corrosion behavior of Ni-based superalloy with higher Cr contents-part II. Mechanism of hot corrosion behavior. J. Univ. Sci. Tech. Beijing 3, 91 (1996) 13. M. Li, High Temperature Corrosion of Metals (Metallurgical Industry Press, Beijing, 2001), p. 263. (in chinese) 14. J. Wang, C. Li, T. Zhang et al., Hot corrosion behavior of Ni-Cr-W based superalloy in molten salt environment. Aerosp. Mater. Technol. 44(6), 26–29 (2014). (in chinese)

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15. J. Sun, The oxidation behavior of Ni-20Cr-Si and Ni-20Cr-Si-Al alloys at 1200 °C. J. Chin. Soc. Corros. Prot. (1), 53–58 (1993), (in chinese) 16. H. Chao, L. Yong, W. Yan et al., Hot corrosion behavior of Ni-xCr-6.8Al based alloys. Trans. Nonferrous Met. Soc. China 21(11), 2348–2357 (2011)

Effects of Extrusion Conditions on Microstructure and Age-Hardening Behaviors of Al–Zn–Mg Alloy Y. L. Wang, H. C. Jiang, D. Zhang and L. J. Rong

Abstract Hot extrusions were performed to study the effects of extrusion conditions on microstructure, age-hardening behaviors and mechanical properties of Al–Zn–Mg alloy. The increasing of extrusion ratio from 10 to 35 effectively enhanced solid solubility of the as-extruded sheets. It accelerated the age-hardening response and significantly increased the number density of nanoscale matrix precipitates in peak-aged sheets. Higher extrusion ratio led to narrower extruded fibers in the sheets, while the extrusion temperature showed the opposite effect. The number of //ED grains significantly decreased while that of //ED grains increased a lot with increasing of extrusion ratio. Due to the differences in degree of deformation, grain size decreased monotonously from center to surface, thus contributing to the higher strength near the surface of peak-aged sheets than the center. The mechanical properties can be more homogeneous on cross section of the sheets by appropriately increasing extrusion ratio and temperature.

1 Introduction Among multifarious light-weight materials, aluminum alloys exhibit excellent workability and extrusion process is a common shape-forming procedure for aluminum products [1]. 7xxx series aluminum alloys (Al–Zn–Mg–(Cu) alloys) are widely used as extruded profiles in rail transit, automobile, and aerospace industry due to their fast age-hardening response [2–4]. Several published works reported that extrusion conditions can significantly influence the microstructure of extruded aluminum alloys and further affect their mechanical properties [1, 5–7]. However, we found that the age-hardening response of Y. L. Wang · H. C. Jiang (B) · D. Zhang · L. J. Rong CAS Key Laboratory of Nuclear Materials and Safety Assessment, Institute of Metal Research, Chinese Academy of Science, Shenyang 110016, China e-mail: [email protected] Y. L. Wang · D. Zhang School of Materials Science and Engineering, University of Science and Technology of China, Hefei 230026, China © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_17

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Al–Zn–Mg alloy can also be markedly impacted by extrusion conditions. Unfortunately, related works always concentrate on the grain features but ignore the difference of precipitating behaviors after extrusions in different conditions. At the same time, mechanical properties can be variable on different positions of extruded profiles. Therefore, the extruded profiles with more homogeneous mechanical properties on their cross section are much better for practical applications. This paper is aimed at revealing the relationship between extrusion conditions and microstructure as well as age-hardening behaviors of Al–Zn–Mg alloy and improving the hot extrusions to obtain extrudates with more homogeneous mechanical properties.

2 Experimental Procedure 2.1 Material and Processing Conditions An Al–Zn–Mg alloy (7N01 aluminum alloy) was used in the current study. The ingots were prepared through permanent mould casting and then were treated by homogenization at 743 K for 24 h. After homogenization, the ingots were hot-extruded into sheets in different extrusion conditions including three extrusion ratios and two extrusion temperatures. The detailed extrusion conditions were given in Table 1. Then the extruded sheets were immediately quenched with water at room temperature, followed by 72 h natural aging to stabilize their microstructure. The subsequent artificial aging was carried on the two-stage aging regime, i.e., 373 K/12 h + 423 K/x h.

2.2 Experimental Methods To analyze the grain morphology of extruded sheets, specimens were systematically observed on the OLYMPUS GX51 optical microscope (OM) from different directions. The transmission electron microscope (TEM), JEM 2100F, was used to study the nanoscale precipitates that can form during artificial aging. The crystallographic

Table 1 Detailed extrusion conditions of hot extrusion

Conditions

Extrusion ratio

Extrusion temperature (K)

Ram speed (mm/s)

R10-T723

10

723

14

R22-T723

22

723

14

R35-T723

35

723

14

R35-T753

35

753

14

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data of extruded sheets was collected by the ZEISS MERLIN COMPACT scanning electron microscope (SEM) through the electron backscattering pattern (EBSP) method. Hardness test was carried on the Micromet 5103 micro-Vickers tester to investigate the age-hardening behaviors. The specimens were sampled from the center of extruded sheets. For each point of aging time, a load of 0.5 Kgf and a hold time of 15 s were applied and five locations on the longitudinal section of each specimen were measured to get average value. Tensile test was conducted on the electronic universal testing machine with a strain rate of 2 mm/min. Specimens for the tensile test were sampled along the longitudinal direction of experimental sheets. Differential scanning calorimetry (DSC) test was performed on the Netzsch STA 449F3 thermal analyzer to study the precipitation process. Specimens for DSC test were polished to disks of 3 mm in diameter and 1 mm in thickness, and they were heated up from 350 to 650 K with a heating rate at 10 K/ min.

3 Results 3.1 Age-Hardening Behaviors The time-dependent age-hardening behaviors of the second-stage aging are characterized by the hardness test, as shown in Fig. 1. Every age-hardening curve shows two hardness peaks, which is similar to the result reported by the past studies [4, 8]. Significantly, the corresponding time for hardness to reach the first peak value decreases from 10 to 6 h when the extrusion ratio varies from 10 to 35. It indicates that the higher the extrusion ratio, the faster the age-hardening response will be. Meanwhile, hardness of specimen rises with increasing of the extrusion ratio. However, the age-hardening behaviors of specimens at two extrusion temperatures are quite synchronous, which means that extrusion temperature shows no obvious effect on age-hardening response. Additionally, age-hardening curve moves to a negative direction when the extrusion temperature changes from 723 to 753 K.

3.2 Microstructure Figure 2 shows the three-dimensional (3D) OM images of specimens sampled from the center of as-extruded sheets in different extrusion conditions. It can be clearly seen from Fig. 2a–c that the size of grains decreases significantly with increasing of extrusion ratio. As shown in Fig. 2d, severe recrystallization presents when extrusion temperature rises to 753 K. Meanwhile, longitudinal sections of specimens in four extrusion conditions all consist of extruded fibers of which shapes are elongated. Figure 3 shows the statistical result of the width of extruded fiber in specimens.

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Fig. 1 Age-hardening curves of specimens in the second-stage ageing process

Obviously, higher extrusion ratio leads to narrower extruded fibers in the sheets, while extrusion temperature shows the opposite effect. The R35-T723 sheet was employed to study the difference between grain features at different position of extruded sheets. Figure 4 demonstrates the OM images of the typical position on the cross section of the as-extruded R35-T723 sheet. The corresponding grain size given by statistical measurement of each position is shown as inset. It can be easily observed that the grain size decreases from center to surface of the sheet. Grains at the diagonal edge are tinyequiaxed, of which size are only ~3.5 μm, while grains at the center reaches ~16.1 μm in size. Figure 5 illustrates the orientation imaging microscopy (OIM) images of a longitudinal section of as-extruded sheets in different extrusion conditions. The images are shown in the Inverse Pole Figure color. Different colors correspond to different crystal orientations along the extruded direction (ED). It can be obviously observed that all the microstructures of as-extruded sheets exhibit strong texture in different conditions. As demonstrated in Fig. 5a–c, area covered with //ED grains remarkably enlarges while that with //ED grains decreases with the increase of extrusion ratio, which agrees well with the conclusion put forward by Kaneko [1]. Meanwhile, the isolated coarse grains filled with no subgrain are related to abnormal growth of the recrystallized grains derived from hot extrusion. Obviously, as extrusion temperature rises from 723 to 753 K (Fig. 5c and d), abnormal growth of recrystallized grains is promoted significantly. At the same time, by comparing Fig. 5c and e, local coarsening of grains at the surface of the as-extruded sheet is more obvious than that of the center. Aging treatment shows a negligible effect on grain structure but noteworthy impact on precipitates formed during the aging process [4, 9]. Morphology of matrix precipitates (MPts) in peak-aged (first peak) sheets are shown in Fig. 6. As demonstrated in Fig. 6a–e, MPts in center and transverse surface of sheets in different extrusion conditions are similar in size but variable in distribution. MPts in R10-T723 sheet are significantly less in number than other extrusion conditions. In addition, HRTEM image of typical particles in the center of R35-T723 sheet is shown in Fig. 6(f). As

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Fig. 2 3D OM image of specimens sampled from center of as-extruded sheets in different extrusion conditions: a R10-T723; b R22-T723; c R35-T723; d R35-T753. L: longitudinal (extruded) direction; T: transverse direction; S: short-transverse direction

exhibited, it proves the presence of G.P. zones and η phase in peak-aged state [10]. For purpose of statistical study of those MPts, the particle size and number density of them are shown in Fig. 7. Specifically, particle sizes of MPts in all conditions are close to ~5.5 nm. Nevertheless, the number density of MPts increases dramatically with rising of extrusion ratio, while the effects of extrusion temperature and position are comparatively weak.

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Fig. 3 Relationship between width of extruded fiber and extrusion condition

3.3 Mechanical Properties Figure 8 shows the distribution of hardness on cross section of as-extruded sheets dealt with water quenching. Hardness on a quarter of the cross section was measured as representation to characterize the distribution of hardness on whole cross section. It can be easily observed that the hardness in every condition tends to increase from center of the cross section to the surface. As shown in Fig. 8a, the difference of hardness between center and surface in R10-T723 sheet reaches ~33 HV, which is much larger than the other three conditions. Correspondingly in Fig. 8d, the distribution of hardness in R35-T753 sheet is the most homogeneous one that shows the lowest hardness difference varying from ~90 to 104 HV. Moreover, it can be seen that a fraction of the area with the highest hardness near surface of the sheets tends to increase as extrusion ratio rises. Extrusion condition and sampling position can significantly affect the mechanical properties of experimental sheets. Figure 9 exhibits the longitudinal-direction mechanical properties of peak-aged sheets. When sampling position is near center of experimental sheets, both the ultimate tensile strength (UTS) and yield strength (YS) of specimens monotonously increase with the extrusion ratio rising from 10 to 35. Meanwhile, the improvement of strength leads to ductility loss in R35-T723 sheet which shows the lowest elongation (~15.3%) with the extrusion temperature at 723 K. On the contrary, with extrusion temperature increases from 723 to 753 K, UTS and YS decrease while elongation rises to ~16.7%. For the R35-T723 sheet, UTS and YS of sheet surface respectively reach ~425.4 and 375.2 MPa, which is much higher than that of the center. It agrees well with the variation of hardness shown in Fig. 8. Figure 10 shows SEM images of the tensile fracture surface. All the specimens exhibit transgranular fracture and no obvious intergranular cracking feature, proving that ductile fracture happens in all conditions during tensile test. As shown in Fig. 10a,

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Fig. 4 Microstructure evolution on cross section of as-extruded R35-T723 sheet

b and d, the fracture surfaces of center of R10-T723, R22-T723, and R35-T753 sheets consist of large quantity of dimples, which are visibly more in number than that of R35-T723 sheet shown in Fig. 10c. Moreover, the fracture surface becomes more flat than others when the specimen is sampled from surface of R35-T723 sheet, as demonstrated in Fig. 10e.

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Fig. 5 OIM images of longitudinal section of as-extruded sheets in different extrusion conditions: a center of R10-T723 sheet (simplified as C, R10-T723); b C, R22-T723; c C, R35-T723; d C, R35T753; e surface of R35-T723 sheet (simplified as S, R35-T723). f unit triangle for crystallographic orientations

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Fig. 6 TEM images of MPts in different positions of peak-aged sheets in different extrusion conditions: a C, R10-T723; b C, R22-T723; c C, R35-T723; d C, R35-T753; e S, R35-T723. HRTEM image of typical particles in C, R35-T723 is shown in (f)

4 Discussion With the increase of extrusion ratio, the deformation of experimental sheets becomes more severe. Therefore, the width of extruded fibers on the longitudinal direction of extruded sheets decreases significantly. According to the Hall–Petch relation, the strength is inversely proportional to the square root of the grain size. For the tensile

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Fig. 7 Statistical result of particle size and number density of MPts in different positions of peakaged sheets in different extrusion conditions: a particle size; b number density

Fig. 8 Distribution of hardness on cross section of as-extruded sheets in different extrusion conditions: a R10-T723; b R22-T723; c R35-T723; d R35-T753

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Fig. 9 Mechanical properties of specimens in different positions of peak-aged sheets in different extrusion conditions

stress along the longitudinal direction, the effective grain size is approximately equal to the width of fibrous grain [11]. DSC analysis was performed on the first-stage aged sheets to further characterize the difference of aging precipitation behaviors in the second-stage aging. Figure 11 shows the DSC curves of the first-stage aged sheets in different extrusion conditions. For all curves, the first endothermic process between 350 and 450 K is associated with formation and coarsening together with dissolution of G.P. zones. After that, the rise of the curves represents the exothermic transformation from G.P. zones to η phases. Then the increase of slope is most likely related to the process that η phases transform to η phases [4]. The overlap for exothermic process for precipitation of η phases and η phases are caused by rapid heating, which leads to the direct transition from η phases to η phases [12]. Interestingly, it can be obviously observed from Fig. 11 that formation of G.P. zones, η phases and η phases all happens ahead of time as extrusion ratio increases, while extrusion temperature shows ignorable influence. It agrees perfectly with the phenomenon shown in age-hardening curves that increase of extrusion ratio significantly accelerates aging response of specimens. In fact, crosssectional size of extruded sheet with higher extrusion ratio is much smaller. Crosssectional size directly affects the cooling process from surface to center of extruded sheet. The smaller the size, the faster it can be cooled. As a result, the extruded sheet with higher extrusion ratio was quickly quenched to room temperature by water and owns higher solid solubility. Thus, it finally leads to faster age-hardening response and provides the extruded sheet with more nanoscale MPts in peak-aged state as mentioned previously in age-hardening curve and the TEM images. Such refining of extruded fibers and high number density of MPts contribute to increase the strength as extrusion ratio rises. However, an increase of extrusion ratio also causes an increase in a number of grain boundaries (GBs). When dislocations glide to MPts or GBs, they can be pinned and then lead to dislocation pile-up. The dislocation pile-up can promote stress concentration and further accelerate the initiation and propagation of cracks under stress. Therefore, elongation tends to drop in experimental sheets with higher extrusion ratio. For the sheets with extruded at 723 and 753 K respectively, the

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Fig. 10 SEM images of fracture surfaces sampled from different positions of peak-aged sheets in different extrusion conditions: a C, R10-T723; b C, R22-T723; c C, R35-T723; d C, R35-T753; e S, R35-T723

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Fig. 11 DSC curves of the first-stage aged sheets in different extrusion conditions

same extrusion ratio leading to similar solid solubility after quenching. Furthermore, it leads to almost synchronous aging response, which means that the time to reach the peak-aged state is the same in the second-stage aging. Thus, both the size and number density of MPts in peak-aged state of the two sheets show almost no difference. Raise of extrusion temperature promotes recrystallization in the experimental sheet through accelerating abnormal growth of recrystallized grains. Such obvious coarsening of grains leads to severe duplex-grain structure and broadens the width of extruded fibers in the sheet extruded at 753 K. It finally makes the strength lower than that of the sheet extruded at 723 K. The sampling position can also significantly affect the microstructure and properties of experimental sheets. First, deformation of surface of extruded sheet during hot extrusion is much severe than that of center. It explains why the measured grain size tends to decrease monotonously from center to surface of R35-T723 sheet. Hence, stronger refine of grains contributes to the higher strength in the specimen sampled from surface. Meanwhile, high density of GBs will increase the resistance for dislocation glide, which is harmful for ductility of surficial sample. The hardness on cross section of as-extruded sheets also varies from different positions. For the four extrusion conditions, the cross-sectional hardness all tends to increase from center to surface of the as-extruded sheets. It coincides with the difference of strength measured from different positions in peak-aged sheets. The difference of grain size on cross section is still the significant factor that dominates the variation of hardness. First, extrusions with lower ratios may enlarge the difference of deformation degree between center and surface of sheets compared to higher ratios. Meanwhile, the quenching process for the sheet with lower extrusion ratio will be much slower than other conditions. It is a benefit for growing of interior grains and can enlarge the difference of grain size between surface and center of the sheets. Therefore, the sheet with the lowest extrusion ratio (10) shows the highest difference of cross-sectional hardness between surface and center, which is more than ~30 HV. Nevertheless, stress concentration can also affect hardness distribution feature on cross section of as-extruded sheets. First, overly inhomogeneous deformation

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during hot extrusion can induce stress concentration to interior of the sheets. At the same time, the quenching can also aggravate the stress concentration in asextruded sheets. However, stress concentration within interior of the sheets can be released by recovery and recrystallization processes, which can be promoted by severe deformation and high temperature. Therefore, distribution of cross-sectional hardness is more homogeneous in the sheet extruded at higher ratio and temperature. It indicates that extrusions with an appropriately higher ratio and temperature are benefited for obtaining more homogeneous mechanical properties on cross section of extrudates.

5 Conclusions (1) Increase of extrusion ratio from 10 to 35 enhances solid solubility of as-extruded sheets. It accelerates the age-hardening response and significantly increases the number density of nanoscale matrix precipitates in peak-aged sheets (2) As extrusion ratio increases, width of extruded fibers drops obviously. Rise of extrusion temperature from 723 to 753 K enlarges fiber width by promoting recrystallization. (3) The microstructures of as-extruded sheets exhibit strong texture in different extrusion conditions. As extrusion ratio increases, number of //ED grains significantly decreases while that of //ED grains increases a lot. (4) The mechanical properties can be more homogeneous on cross section of the sheets by appropriately increasing extrusion ratio and temperature. Acknowledgements The authors acknowledge the financial supports by National Key R&D Program of China (No. 2016YFB1200504) and Strategic Priority Program of the Chinese Academy of Sciences (No. XDB22000000).

References 1. S. Kaneko et al., Effect of the extrusion conditions on microstructure evolution of the extruded Al–Mg–Si–Cu alloy rods. Mater. Sci. Eng. A 500(1–2), 8–15 (2009) 2. J.C. Williams, E.A. Starke Jr., Progress in structural materials for aerospace systems 1. Acta Mater. 51(19), 5775–5799 (2003) 3. J.G. Tang et al., Influence of quench-induced precipitation on aging behavior of Al-Zn-Mg-Cu alloy. T. Nonferrous Met. Soc. 22(6), 1255–1263 (2012) 4. Y.L. Wang et al., Two-stage double peaks ageing and its effect on stress corrosion cracking susceptibility of Al-Zn-Mg alloy. J. Mater. Sci. Technol. 34(7), 1250–1257 (2018) 5. J. Zhou, J. Duszczyk, Effect of extrusion conditions on mechanical properties of Al-20Si-3Cu1Mg alloy prepared from rapidly solidified powder. J. Mater. Sci. 26(14), 3739–3747 (1991) 6. S. Kumar et al., Effect of extrusion parameters on the microstructure and properties of an Al-Li-Mg-Zr alloy. J. Mater. Sci. 29(4), 1067–1074 (1994) 7. S. Karabay et al., Investigation extrusion ratio effect on mechanical behaviour of extruded alloy AA-6063. J. Mater. Process. Technol. 135(1), 101–108 (2003)

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8. C.D. He et al., Precipitation process and effect of precipitated phases of 7050 aluminum alloy two-stage double peak aging. Mater. Mech. Eng. 35(6), 38–41 (2011). (in Chinese) 9. X.Y. Sun et al., Correlations between stress corrosion cracking susceptibility and grain boundary microstructures for an Al-Zn-Mg alloy. Corros. Sci. 77, 103–112 (2013) 10. Y.L. Wang et al., Variation of nanoparticle fraction and compositions in two-stage double peaks aging precipitation of Al−Zn−Mg alloy. Nanoscale Res. Lett. 13(1), 131 (2018) 11. H. Adachi et al., Effect of hot-extrusion conditions on mechanical properties and microstructure of P/M Al-Zn-Mg-Cu alloys containing Zr. Mater. Sci. Forum, 1479–1484 (2006) 12. A. Deschamps et al., Influence of predeformation on ageing in an Al–Zn–Mg alloy—I. Microstructure evolution and mechanical properties. Acta Mater. 47(1), 281–292 (1998)

Process Optimization Design of High-Strength Ag–Cu–Ni Alloy Based on Orthogonal Experiments Helong Hu, Haibin Li, Wenjun Yu, Yongzhen Jiao, Tingyi Dong and Baoguo Lv

Abstract Silver–copper–nickel alloy has been wildly used in the military field as energy transfer component and signal transmission parts for its high conductivity and excellent mechanical properties. Rolling and heat treatments parameters influence the alloy properties effectively. However, systematic study on the parameter influence on the alloy properties and microstructure was seldom conducted. Herein, the influence of deformation, temperature of heat treatments and cooling mode on the alloy properties and microstructure were studied by orthogonal experiments design. The results showed that the order of the effects of processing parameters on the mechanical properties was deformation, annealing temperature, and homogenizing temperature. The hardness was the highest with 155 Vickers hardness, when the alloy was treated with homogenizing treatment temperature of 720 °C, annealing temperature of 250 °C, and cold rolling deformation of 55%. The microstructure observations showed that the alloy presented eutectic structure and distributed along the direction of deformation.

1 Introduction Pure silver was used as electrical contact materials for its excellent performance, such as easy processing, good plasticity, well conductivity and thermal conductivity, excellent oxidation resistance, and so on [1]. Silver and silver alloy were the most wildly used, the most economical and the most basic precious metal-based electrical contact materials, which were wildly used in moderate and heavy loading instruments [2]. Pure silver electrical contact materials have several disadvantages, such as short service life, unreliable operation, due to its low melting point, low hardness, and poor reliability. Thus, silver-based alloy with good abrasion resistance, well fusion H. Hu (B) · H. Li · W. Yu · Y. Jiao · T. Dong · B. Lv Beijing Trillion Metals Co., Ltd., Beijing, China e-mail: [email protected] H. Li · T. Dong · B. Lv GRIKIN Advanced Material Co., Ltd., Beijing, China © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_18

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Levels

Factors Homogenizing Deformation heat treatment (%) temperature (°C)

Annealing temperature (°C)

1

690

45

150

2

720

55

200

3

750

65

250

welding resistance, and excellent arc resistance was desperately in need. In addition, good physical, chemical and processing properties were indispensable [3]. The most fundamental requirements for the electrical contact materials were well electrical conductivity and thermal conductivity, high hardness, and excellent abrasion resistance [4]. In order to improve those performances, copper was added to hardness, strength, and abrasion resistance, nickel was also added to refine grains, improve strength, and abrasion resistance [5]. In actual production, Ag–Cu-based solid solution strengthening materials and Ag–Ni-based dispersion strengthening materials formed [6]. Rolling and heat treatment parameters affect the performance of silver–copper–nickel alloy obviously [7]. But, the documents related to the influence of rolling and heat treatments parameters on the properties and microstructure of silver— copper–nickel alloy were insufficient. In this study, the influences of deformation, homogenizing treatments temperature, and annealing treatments on the properties and microstructure of alloy were investigated based on orthogonal experiments.

2 Experiments The raw materials were industrial pure silver, electrolytic nickel, and electrolytic copper. The ration was Ag:Cu:Ni  78%:20%:2%. The raw materials were melted in vacuum induction furnace and pouring into the copper mould, and alloy ingot was obtained. The ingot was homogenized in the muffle furnace. Samples with a size of 50 * 50 * 10 mm were cut off from the alloy ingot. The samples were treated with different deformation, heat treatments temperature. Table 1 shows designed orthogonal factor level table of three factors and three levels. Based on the phase diagram, the alloy melting point was about 790 °C, thus three temperature points 690, 720, 750 °C were selected. Hardness test was carried on microhardness tester. CSS-44100 electronic universal testing machine was adopted to test the mechanical properties with a stretching rate of 2 mm/min. The sample prepared was based on GB/T228-2002. Microstructure of the alloy was carried on the Leica metallographic microscope.

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Table 2 Orthogonal test results Serial number

Homogenizing Deformation Annealing temperature (%) temperature (°C) (°C)

Tensile strength (MPa)

Elongation Microhardness (%) (HV)

1

690

45

150

380

15

120

2

690

55

200

415

18

130

3

690

65

250

408

24

127

4

720

45

200

455

23

141

5

720

55

250

490

29

155

6

720

65

150

480

35

150

7

750

45

250

350

12

109

8

750

55

150

371

16

115

9

750

65

200

379

19

119

3 Results and Discussion 3.1 Orthogonal Test and Range Analysis Orthogonal test were established, results and range analysis results were shown in Tables 2 and 3. The results showed that effects of processing parameters to the tensile strength, elongation and hardness. The order of the effects of processing parameters to the mechanical properties was deformation, annealing temperature and homogenizing temperature.

3.2 Effects of Processing Parameters to the Alloy Mechanical Properties 3.2.1

Effects of Homogenizing Temperature to the Alloy Mechanical Properties

The effects of homogenizing temperature to the alloy mechanical properties were showed in the tendency chart, based on the orthogonal test. It is showed that the tensile properties, elongation and microhardness increase at first and then decreased. The tensile properties, elongation and microhardness reached 475 MPa, 29%, 149 HV at the temperature of 720 °C (Fig. 1). Figure 2 showed the metallograph of the alloy after homogenizing treatment at 750 °C. The nickel precipitated was circled in black. When the temperature of homogenizing treatments is too high, nickel precipitated, the mechanical properties and microhardness decreased, and the ductility deteriorated. When the temperature

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Table 3 Range analyses Tensile strength (MPa)

A

B

C

K1

401

395

410

K2

475

425

416

K3

367

422

416

R

108

30

6

Factor

Order: A > B > C

Elongation

K1

19

17

22

K2

29

21

20

K3

16

26

24

R

13

9

4

Factor

Order: A > B > C

Microhardness (HV)

K1

126

123

128

K2

149

134

130

K3

113

132

131

R

36

11

3

Factor

Order: A > B > C

500

160

475

475

149

145

450 425

130

126

401

400

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of the homogenizing treatments is too low, composition segregation formed in the process of melting could not be homogenized effectively. Thus, the temperature of homogenizing treatments should not be too high.

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The Effects of Deformation to the Alloy Mechanical Properties

Figure 3 showed the effects of the deformation to the alloy mechanical properties. It is showed that the tensile strength, elongation, and microhardness are raised by 30 MPa, 9%, 11 HV. Increasing the deformation can cause dislocation pile-up and work hardening. Thus, the microhardness increased rapidly. The plasticity can be improved through annealing treatments. Figure 4 showed the microstructure of the alloy after rolling. The microstructure results showed that the alloy presented eutectic structure and distributed along the direction of deformation.

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The Effects of Annealing to the Alloy Mechanical Properties

Figure 5 showed the effects of annealing to the alloy mechanical properties. It can be concluded from the figure that the alloy microhardness increased as the annealing

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temperature. This can be attributed to that the microhardness increased as the annealing temperature increased at low temperature annealing. This phenomenon was verified by experiments, but the mechanism was unknown. The alloy microhardness reached the highest level of 155 HV after annealing at the temperature of 250 °C.

3.3 Optimal Solution The experiment results showed that the best alloy mechanical properties listed as this: tensile strength was 490 MPa, elongation was 35%, microhardness was 155 HV. In consideration of comprehensive mechanical properties and feasibility of processing parameters, the best technical parameters are listed as following: the homogenizing temperature was 720 °C, the deformation was 55%, and the annealing temperature was 250 °C.

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4 Conclusion 1. The order of the effects of processing parameters to the mechanical properties is deformation, annealing temperature, and homogenizing temperature. 2. The best technical parameters are listed as following: the homogenizing temperature is 720 °C, deformation is 55%, and annealing temperature is 250 °C. Under those processing parameter, the tensile properties, elongation and microhardness are 490 MPa, 35%, 155 HV.

References 1. R.W. Lampe, in Compensation of Phase Errors due to Coaxial cable Flexure in Near-Field Measurements. 1990 IEEE Antennas propagate Society of International Symposium Digest (Dallas, TX: IEEE Press, 1990), pp. 1332–1324 2. A.C. Newell, Error analysis techniques for planar near-field measurement. IEEE Trans. Antennas Propag. 36(6), 754–768 (1988) 3. C.D. Des Forges, Sintered materials for electrical contacts. Powder Met. 22(3), 138–143 (1979) 4. R. Michal, K.E. Saeger, Metallurgical aspects of silverbased contact materials for air-break switching devices for power engineering. IEEE Trans. 12(1), 71–75 (1989) 5. S. Kang, C. Brecher, P. Womgert, Erosion resistant Ag-SnO2 electrical contact material: U.S. 4904317[P]. 1990-02-27 6. P.B. Joshi, Improved pm silver-zinc oxide electric calcontacts. Met. Powder Rep. 54(3), 37–41 (1999) 7. A.L. Stroyuk, V.V. Shvalagin, S.Y. Kuchmi, Photochemical synthesis, spectral-optical and electrophysical properties of composition nanoparticles of ZnO/Ag. Theor. Exp. Chem. 40(2), 982101 (2004)

Strain Distribution and Metal Flow of Bulk Forming of Molybdenum Yihang Yang, Ailong Zheng, Zhimin Huang, Fusheng Peng and Houan Zhang

Abstract The characteristics of bulk-forming of refractory metal were analyzed, and the effects of strain paths and microstructures on the Mo bulk-forming were discussed. By means of finite element method, metallographic and Electron Backscattered Diffraction (EBSD), the flow and strain distribution of the Mo bulk after multidirection forging were analyzed. It was found that the metal flow in the core of bulk was intense, and the strain storage energy was the largest. The grain sizes in the core were also larger than those around edges, and there were still some deformed textures in the material. The metallographic results were similar to the Finite Element Method (FEM) results.

1 Introduction Refractory metals are widely used in the fields of aerospace, nuclear industry, national defense, and semiconductors because of their high melting points, high density, good heat conduction, and electrical conductivity. With the development of engineering technology, refractory metals with fine microstructure, large-sized scale, and excellent mechanical properties are becoming more and more important. However, traditional manufacturing processes were accompanied by low density, coarse grains, or small product scale, which hampered their development seriously in these fields.

Y. Yang · H. Zhang (B) Fujian Key Laboratory of Functional Materials and Applications, School of Materials Science and Engineering, Xiamen University of Technology, Xiamen 361024, China e-mail: [email protected] Y. Yang · A. Zheng · Z. Huang · F. Peng Department of R&D, Xiamen Honglu Tungsten Molybdenum Industry Co. Ltd., Xiamen 361021, China H. Zhang Fujian Collaborative Innovation Center for R&D of Coach and Special Vehicle, Xiamen University of Technology, Xiamen 361024, China © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_19

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Recently, bulk fine-grained materials processed by large plastic deformation have attracted interests by many metal engineers and technicians. According to their opinions, multidirectional forging seems to be a great potential for fabricating blocks that can be suitable for industrial applications [1, 2]. Previous researches about multidirectional forging mainly focused on Mg alloys [3, 4], Fe alloys [5–7], Al alloys [8–10] and Ti alloys [11]. However, little research has been conducted in the field of manufacturing refractory metals by multidirectional forging. In this work, the multidirectional forging has been performed on the Mo bulk material. Strain distribution and metal flow of plastic-deformated Mo were discussed in details.

2 Experimental Procedure The raw material used for this study was supplied by Xiamen Tungsten Co., Ltd., China. Results from chemical composition tests of raw material are shown in Table 1. The content of Mo is more than 999.5‰. The schematic diagram is shown in Fig. 1 represents to complete cycle of multidirectional forging process [12]. During the process, the Mo bulk was rotated by 90° after each pass. The bulk was first homogenized at a temperature of 1100 °C for 30 min in the atmosphere of H2 . Multidirectional forging was performed in a 100 kN forging press. The bulk was held in the furnace at the temperature of 1100 °C for 10 min after  each pass. Equivalent strain in one pass ε  −0.5, Cumulative strain in one cycle ε  |ε1 + ε2 + ε3 |  −1.5. The multidirectional forging was performed at a strain rate of about 10 s−1 and an equivalent total true strain of 300% (0.5 × 6; 6 passes for 2 cycles) was obtained for each bulk. After the completion of the forging operation (2 cycles  6 passes) each forged bulk was cooled in air. The microstructural analysis was performed by optical microscope (Keyence VHX500F), SEM and EBSD. EBSD analysis was performed using Oxford instruments Symmetry.

Table 1 Chemical composition (wt.‰) of the pure Mo bulk Element

Al

Si

Ni

Cr

C

Fe

Ti

O

Ca

Mg

N

W

wt.‰

0.018 0.024 0.022 0.022 0.026 0.026 0.022 0.024 0.016 0.016 0.016 0.13

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Fig. 1 Schematic illustration of multidirectional forging for 1 cycle

3 Results and Discussion 3.1 FEM Analysis First of all, FEM methodology was applied to find out the materials flow during the multidirectional forging process. 3D models were constructed by Inventor 3D software, and they were including pressure head and Mo bulk for every stage. Then STL file was imported to DEFORM-3D software to preset the process and simulate the analysis. Figure 2 shows the distribution of strain-effective and grain model after plastic deformation with true strain of 300%. The strain-effective in the core of forged Mo was about 5, compared with about 1 in the edges.

Fig. 2 FEM of multidirectional forged Mo bulk

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Fig. 3 Extrinsic feature of the processed molybdenum

3.2 Appearance Figure 3a illustrates the bulk photo of molybdenum by two cycles processed multidirectional forging. This bulk was 80 mm in diameter, 50 mm in height. Almost no cracks observed on the face of the bulk. Mo products were fabricated after further processing, as shown in Fig. 3b. The density obtained by METTLER TOLEDO is shown in Table 2. To analyze the density of the specimen, they were cut with size of 5 × 10 × 10 mm. The Mo bulk processed by multidirectional forging attains 99.7% of theoretical density, about 10.17 g/cm3 . It showed more optimal results than the processes of rolling [13], upset [14], extrusion, etc.

3.3 Mechanical Properties To evaluate the tensile properties of the Mo bulk, tensile test/tested specimens were shown in Fig. 4a. These specimens were tested in tension using the MTS instrument. Tests were carried out at room temperature with a loading rate of 1 × 10−3 m/min. Figure 4 shows that the tensile fracture partially occurred, as evidenced by some apparent necking (marked with red circle). It can be observed that the room

Table 2 Density of molybdenum before and after processing (g/cm3 ) No. 1 Before processing Multidirectional forged

No. 2

No. 3

Mean value

9.875

9.864

9.871

9.870

10.168

10.173

10.166

10.169

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Fig. 4 Tensile test specimens (a) and summary of the strengths (b)

temperature tensile strength was about 640 MPa, and the elongation was about 25%, which indicated that the Mo bulk has a good plasticity at room temperature.

3.4 Microstructural Investigation Optical microscope was conducted to study the evolution of the particles after multidirectional forging. Samples for microstructure test were carried out from the central part of bulk to the edge. “a” track was parallel to the last compression axis, “b” track was a’s vertical direction, as shown in Fig. 5. A pure Mo bar with ϕ80 mm in diameter was prepared by sintering. Samples were cut into size 8 × 10 × 10 mm from the center of the ingot. A certain amount of porosity appeared in Mo powders sintering. The grain size of equiaxed grain was 20–40 µm, as shown in Fig. 6. The “a” & “b” directional metallographic phase of multidirectional forged Mo bulk were analyzed, as shown in Figs. 7 and 8. Near-equiaxed grains were presented in the microstructure. It was found that the grain size of the core is large, compared with smaller around the edge. This was corresponding to the FEM results. Furthermore, the metal flow in the core of multi-direction forging was intense, and the strain storage energy was the largest.

Fig. 5 Identification of the direction of metallographic phase

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Fig. 6 Metallographic phase of raw material

EBSD analysis has been carried out for the multidirectional forged Mo sample, which also showed distinctly refined grains with size of 2–10 µm, as shown in Fig. 9. A large number of small-angle grain boundaries exist in bigger grains. Because of its large accumulative strain (300%), the storage energy of dislocation and the driving force of recrystallization were enhanced. When the sintered coarse Mo grains were repeatedly multidirectional forged, they finally recovered and recrystallized to generate fine equiaxed grains. The colors are determined by the orientation of each grain, which shows there were some deformed textures in the material. Therefore, texture cannot be eliminated by multidirectional forging in the field of refractory metals.

4 Conclusions In this study, we have investigated the strain distribution and metal flow of Mo bulk after being multidirectional forged for two cycles. On the basis of the results and their analysis, the following conclusions can be ascertained: (1) A bulk with 80 mm in diameter, 50 mm in thickness has been made by multidirectional forging. The Mo bulk attains 99.7% of theoretical density. The MDFed Mo bulk presents fine equiaxed grains in microstructure. (2) The Mo grain size is significantly refined after accumulative strain reaches 300%, resulting in a uniform fine-grained structure with grain size of 2–10 µm. The excellent tensile properties are achieved after six passes forging, with

Fig. 7 Metallographic phase of forged Mo bulk (a direction in Fig. 5)

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Fig. 8 Metallographic phase of forged Mo bulk (b direction in Fig. 5)

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elongation to fracture 25% and ultimate tensile strength 640 MPa at room temperature. (3) The grain sizes of the core are larger than that around edges, and there are still some deformed textures in the material. Material flow in the core of multidirection forged Mo bulk is intense, and the strain storage energy is the largest. The metallographic results are similar to the FEM results. Acknowledgements This work was supported by the National Key R&D Program of China (2017YFB0305603), the Education and Research Project of Fujian Province for Youths (JAT170401), the Program for Innovative Research Team in Science and Technology in Fujian Province University. We thank senior engineer Zheng Ailong for fruitful discussions, Luo Xinhua for technical support.

References 1. O. Sitdikov, T. Sakai, A. Goloborodko et al., Grain fragmentation in a coarse-grained 7475 Al alloy during hot deformation. Scr. Mater. 51, 175–179 (2004) 2. O. Sitdikov, A. Goloborodko, T. Sakai et al., Grain refinement in as-cast 7475 Al alloy under hot multiaxial deformation. Mater. Sci. Forum, 381–386 (2003) 3. Q. Chen, D. Shu, C. Hu et al., Grain refinement in an as-cast AZ61 magnesium alloy processed by multi-axial forging under the multi temperature processing procedure. Mat. Sci. Eng. A 541, 98–104 (2012) 4. X.S. Xia, M. Chen, L.U. Yong-Jin et al., Microstructure and mechanical properties of isothermal multi-axial forging formed AZ61 Mg alloy. T. Nonferros Met. Soc. 23, 3186–3192 (2013) 5. B.J. Han, Ultra-fine grained Fe-32%Ni alloy processed by multi-axial forging. Adv. Mater. Res. 97–101, 187–190 (2010) 6. B. Han, X. Zhou, Microstructural evolution of Fe–32%Ni alloy during large strain multi-axial forging. Mat. Sci. Eng. A 447, 119–124 (2007) 7. B. Han, Z. Xu, Grain refinement under multi-axial forging in Fe–32%Ni alloy. J. Alloy. Compd. 457, 279–285 (2008) 8. R. Kapoor, A. Sarkar, R. Yogi et al., Softening of Al during multi-axial forging in a channel die. Mat. Sci. Eng. A 560, 404–412 (2013)

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9. M. Montazeri-Pour, M.H. Parsa, H.R. Jafarian et al., Microstructural and mechanical properties of AA1100 aluminum processed by multi-axial incremental forging and shearing. Mat. Sci. Eng. A 639, 705–716 (2015) 10. M. Montazeri-Pour, M.H. Parsa, Constitutive analysis of tensile deformation behavior for AA1100 aluminum subjected to multi-axial incremental forging and shearing. Mech. Mater. 94, 117–131 (2016) 11. B.F. Ghanbari, H. Arabi, S.M. Abbasi et al., Manufacturing of nanostructured Ti-6Al-4V alloy via closed-die isothermal multi-axial-temperature forging: microstructure and mechanical properties. Int. J. Adv. Manuf. Tech. 87, 755–763 (2016) 12. S. Ghosh, A.K. Singh, S. Mula, Effect of critical temperatures on microstructures and mechanical properties of Nb–Ti stabilized IF steel processed by multiaxial forging. Mater. Des. 100, 47–57 (2016) 13. T. Fujii, Effects of heating rates on primary recrystallization textures in pure molybdenum and TZM alloy sheets. J. Jpn. Soc. Powder Powder Metall. 48, 824–829 (2009) 14. C. Chen, M.P. Wang, P. Jin et al., Effect of annealing temperature on microstructure and transverse ductility of upset pure Mo bars. Chin. J. Nonferrous Met. 19, 1061–1067 (2009)

Cold Deformation Behavior and Mechanical Properties of Forged Pure Nickel N6 Zexi Gao, Zhi Jia, Jinjin Ji, Dexue Liu and Yutian Ding

Abstract With 0.018% content in the crust, pure nickel N6 has excellent conductivity, hot conductivity, and corrosion resistance. At present, it has been widely used in chemical, nuclear, marine, and aerospace fields. However, there are few studies on deformation behaviors and mechanical properties of pure nickel N6 at room temperature. In this paper, cold deformation behaviors and mechanical properties of N6 were studied by tensile, compressive, and impact load experiment at room temperature. With the increase of strain rate, it was found that the work hardening index increases gradually and the strength coefficient showed an upward trend, which basically followed the law of parabolic hardening, and its constitutive equation was obtained after the modification of strain rate. Observation of fracture by scanning electron microscope (SEM), it was found that the tensile fracture way of N6 changed from brittle fracture to ductile fracture and the impact fracture mechanism was both ductile fracture and brittle fracture. This study may lay a foundation for the cold working process of pure nickel N6.

1 Introduction N6 is a pure nickel material with more than 99% content of nickel element. Ni is scheduled for No. 28 element in Periodic Table, belongs to VIIIA Tribe, the third electron layer constituting the atom is basically filled, therefore it can dissolve more alloying elements to be alloyed to maintain the stability of the austenite phase [1–4]. In addition, the nickel element has a face-centered cubic structure, and allotropic Z. Jia (B) · D. Liu · Y. Ding State Key Laboratory of Advanced Processing and Recycling of Nonferrous Metals, Lanzhou University of Technology, Lanzhou, Gansu Province, China e-mail: [email protected] Z. Gao · Z. Jia School of Material Science and Engineering, Lanzhou University of Technology, Lanzhou, Gansu, China J. Ji School of Materials Engineering, Lanzhou Institute of Technology, Lanzhou, Gansu, China © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_20

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transitions are so hard to occur from low temperature to high temperature, and the structure is very stable. The pure nickel N6 is widely applied in coal chemical industry, electronic components, mechanical materials, nuclear power material, and aviation, etc., as high electrical conductivity, hot conductivity, corrosion resistance, and excellent mechanical properties and the plastic processing properties. In the current actual processing and production, the pure nickel N6 are mostly used for the forming of wire rods and bars, however, the formation of wire and bar depends mainly on the cold deformation process of the materials. The cold deformation of the material refers to the plastic deformation behavior of the material occurring under the recrystallization temperature. The plasticity index of a material is usually expressed as the amount of plastic deformation when the material begins to break. It can often be measured by compression test, tensile test, and impact test. In order to avoid sudden brittle fracture, materials with relatively good plasticity can produce proper materials before fracture occurs. By reviewing the current home and abroad literatures, it has been found that research on nickel mainly focuses on hot deformation behavior and hot processing maps. Such as Mo et al. have researched on the hot deformation behavior and tube extrusion for Nickel [1]. The hot deformation behavior and hot processing diagram of as-cast pure nickel N6 was studied by Ji and Jia [2]. Zhang et al. [3] studied the deformation behavior and processing map of flat hot compression of pure nickel N6. Zhu et al. researched in analysis on microstructure and tensile fracture morphology of pure nickel N6 coldrolled sheet [4]. Characterization of the Hot Deformation Behavior of a Newly Developed Nickel-Based Superalloy was studied by Shi et al. [5]. Fatigue crack growth behavior of an alternative single crystal nickel base superalloy was investigated by Palmert et al. [6]. Tang et al. [7] studied the Formation and Evolution of Shear Bands in Plane Strain Compressed Nickel-Based Superalloy. Dynamic Recrystallization Mechanism of Inconel 690 Superalloy during Hot deformation at High Strain Rate was investigated by Bin et al. [8]. Constitutive Model of Supper-Alloy IN625 Based on Extrusion Test was studied by Wang et al. [9]. Hot deformation behavior and processing map of a γ  -hardened nickel-based superalloy was investigated by Zhang et al. [10]. Hot die forging process optimization of superalloy IN718 turbine disc using processing map and finite element method was investigated by Zhang et al. [11]. However, there are few studies on the deformation behavior and mechanical properties of pure nickel N6 at room temperature. In this paper, through room temperature compression, tension and impact loading tests by forging pure nickel N6, its deformation behavior and mechanical properties at room temperature were studied, expect to control cold processing technology during material deformation to get the best performance products.

2 Material and Experimental The used material was a forging N6 with showing the chemical compositions in Table 1.

Cold Deformation Behavior and Mechanical Properties … Table 1 The chemical composition of pure nickel N6 (wt%)

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The original microstructure is shown in Fig. 1, and the different shades exhibited by the metallographic etchant staining, which consists of a single microstructure, reflect the orientation of the different grains to some extent. In order to study the room temperature mechanical properties and deformation behavior of N6, room temperature compression, tensile, and impact tests were performed. The compression test uses 8 * 12 mm cylindrical specimens and compresses at strain rates of 0.1 and 0.01 s−1 respectively. Tensile specimens and impact specimens were prepared according to the requirements of the national standards.

3 Cold Deformation Behavior of Pure Nickel N6 3.1 Work Hardening Rule of N6 The true stress-true strain curve of forged pure nickel N6 is shown in Fig. 2 at different strain rates with the cold deformation of 50%. With the increase of strain, the stress rapidly rises to the vicinity of the yield stress value, and then gradually rises to the steady state, which is consistent with the rheological characteristics of the low stacking fault energy metal. The curve can be roughly divided into a linear hardening stage and a parabolic hardening stage, and follows the work hardening rule of metal plastic deformation. When ε˙  0.1 s−1 , ε  0.7, the deformation resistance is as high as 629 MPa. When the deformation is within the range of 0–50%, the two

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true stress–true strain curves with different strain rates basically coincide. Due to errors in the experimental process (such as the poor lubrication conditions resulting in the uneven deformation, etc.), there are some non-coincidence areas on the curve. Therefore, when cold plastic deformation of N6 occurs at room temperature, the stress changes have a great influence on the alloy, and the strain rate change has almost no effect. Due to the small strain in the linear hardening stage, the parabolic hardening stage used to describe the entire hardening phase of N6. The Hollomon equation σ  Kεn is used to represent the true stress–true strain relationship in the parabolic hardening stage. The Origin data processing software is used to fit the curves, it is found that the curve basically coincides at low strain, while the curve at high strain has deviations. At strain rates of 0.01 and 0.1 s−1 , the fitted intensity factor and work hardening indices were 684.37, 0.2702 and 702.05, 0.25467, and the correlation coefficients were 0.99462 and 0.99089.

3.2 Constitutive Equation of Pure Nickel N6 In the actual processing, taking into account the effect of strain rate on the work hardening of the material, the exponential hardening equation is modified to σ  K ε(n 1 + n 2 lnε) ε˙ m , and take the logarithm of the equations on both sides and simplify it to

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lnσ  lnK + n 1 lnε + n 2 (lnε)2 + mln˙ε

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(2)

The true stress corresponding to the true strain is calculated from the cold deformation constitutive equation of N6. The comparison between the analog value and the experimental value is shown in Fig. 4. The comparison results show that the constitutive equation has universal applicability.

3.3 Calculation of Impact Toughness Impact toughness is the basic index for measuring the ability of material to resist impact load, that is, the impact energy (Ak) is measured when the impact load sample is broken. The practical significance of the index of impact toughness lies in revealing the brittle tendency of the material, and reflecting the ability of metal material to resist additional impact load. The experimental data for this test is shown in the following Table 2. As shown in Table 2, it can be found that the impact toughness values of the three specimens are not much different and they basically fluctuate within a range.

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Fig. 4 Comparison of analog value and experimental value of true stress-true strain curve for N6 Table 2 Impact toughness calculation table Test number

Meter reading (kgf/m)

Impact distance (m)

Impact work (J)

Impact toughness (J/cm2 )

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132.9

1.63

15.974

19.9675

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132.6

1.69

16.562

20.7025

3

132.0

1.82

17.836

22.295

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Obviously, the impact toughness of three specimens is basically the same. The impact toughness of pure nickel is large, so it belongs to ductile material.

4 Fracture Morphology Observation 4.1 Tensile Specimen Fracture Analysis As shown in Fig. 5, it is the macroscopic morphology of the tensile test fracture.

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(a)

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(a) Stretch rate is 0.0003 m/s

(b) Stretch rate is 0.00008 m/s

(c)

(c) Stretch rate is 0.00001 m/s Fig. 6 Tensile specimen fracture SEM

Obviously, the sample undergoes plastic deformation before it breaks. The initial grain is stretched, broken and caused necking. The surface of the cup-cone fracture surface is rough and gray with no metallic luster fibrous. As shown in Fig. 6, the tensile fracture is a dimple fracture. The dimples in the central fiber region are larger and deeper, however, the dimples of the peripheral shear lip are smaller and shallower. The membrane of dimples has obvious serpentine slip characteristics. With the decrease of the deformation rate, the dimple becomes more and more dense, and the deeper it is, the transition from brittle fracture to ductile fracture, the plasticity of the material becomes better and better.

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Fig. 7 Macroscopic morphology of fracture in impact toughness test

4.2 Impact Toughness Test Fracture Analysis As shown in Fig. 7, it is the macroscopic morphology of the impact toughness test fracture. Macroscopically, there are no major differences in the fracture morphology of the three test specimens. The fractures show a gentle slope, and many silver-white spots are evenly attached to its surface. There are no obvious macroscopic plastic deformation characteristics. As shown in the Fig. 8, the fractures are dominated by intergranular and transcrystalline brittle fractures, some shallow dimples are unevenly distributed on the crystal face, and a large number of tear zones can be observed between the intergranular fractures. There are obvious cleavage stages in the river patterns. As the impact load increases, the dimples become shallower and almost disappear, and more and more tear zones and cleavage steps appear between the intergranular fractures in the crack propagation zone. Ductile fractures gradually turn to brittle fractures.

5 Conclusion (1) Through the room temperature compression test, the strength and work hardening indices of N6 are respectively 684.37, 0.2702, and 702.05, 0.25467 at strain rates of 0.01 and 0.1 s−1 . (2) As the deformation rate decreases, the brittle fracture of the tensile fracture turns into ductile fracture according to the room temperature tensile test. (3) In the impact test, as the impact work increases, the fracture changes from ductile fracture to brittle fracture. Acknowledgements Zhi Jia would like to gratefully acknowledge the support of the National Nature Science Foundation of China (No. 51665032) and Science Foundation for Distinguished Young Scholars of Gansu Province (18JR3RA134).

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Fig. 8 Impact specimen fracture SEM

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References 1. G. Mo, F. Pan Zen, Research on the hot deformation behavior and tube extrusion for Nickel. J. Netshape Form. Eng. (2016) 2. J.J. Ji, Z. Jia, W. Hu, X. Ru, F. Guo, Study on Hot Deformation behavior and hot processing diagram of as cast pure Nickel N6. Hot Work Technol. (2017) 3. B. Zhang, L.L. Zhu, K.S. Wang, W. Wang, X. Ya, F. Hao, The deformation behavior and processing map of flat hot compression of pure nickel N6. Chin. J. Eng. 37(4), 480–487 (2015) 4. L. Zhu, X. Liu, G. Sun, Analysis on microstructure and tensile fracture morphology of pure nickel N6 cold-rolled sheet. Hot Work Technol. 44(5), 247–249 (2015) 5. Z. Shi, X. Yan, C. Duan, C. Tang, E. Pu, Characterization of the hot deformation behavior of a newly developed nickel-based superalloy. J. Mater. Eng. Perform. 27(4), 1763–1776 (2018) 6. F. Palmert, J. Moverare, D. Gustafsson, C. Busse, Fatigue crack growth behavior of an alternative single crystal nickel base superalloy. Int. J. Fatigue, 166–181 (2018) 7. B. Tang, L. Xiang, L. Cheng, D. Liu, J. Li, The formation and evolution of shear bands in plane strain compressed nickel-base superalloy. Metals 8(2), 141 (2018) 8. B. Wang, S.-H. Zhang, M. Cheng, H.-W. Song, Dynamic recrystallization mechanism of Inconel 690 superalloy during hot deformation at high strain rate. J. Mater. Eng. Perform. 22(8), 2382–2388 (2013) 9. Z.T. Wang, S.H. Zhang, M. Cheng, D.F. Li, Constitutive model of supper-alloy IN625 based on extrusion test. Adv. Mater. Res. 314–316, 819–822 (2011) 10. H. Zhang, K. Zhang, Z. Lu, C. Zhao, X. Yang, Hot deformation behavior and processing map of a γ -hardened nickel-based superalloy. Mater. Sci. Eng. A (2014) 11. H.Y. Zhang, S.H. Zhang, Z.X. Li, M.F. Cheng, Hot die forging process optimization of superalloy IN718 turbine disc using processing map and finite element method. Proc. Inst. Mech. Eng. 224(B1), 103–110 (2010)

Microstructure and Properties of Graphene/Copper Matrix Composites Prepared by In Situ Reduction Xu Ran, Yutong Wang and Yong Wang

Abstract Graphene/copper composite powders were prepared by in situ reduction method. The composite powders were mixed with copper powders, and the graphene/copper-based composites were prepared by spark plasma sintering (SPS) process. In this paper, the effects of graphene content on the microstructure and properties were studied. It was found that at certain sintering temperature, with the increase of graphene content, the hardness of the composite material first increased and then decreased. The rates of the relative density and electrical conductivity gradually decreased. When the content of graphene was 1.2 vol%, the composite material had the best comprehensive performance. The microhardness was 116 HV, which was 73% higher than that of pure copper (67 HV). The relative density and electrical conductivity were 98.5 and 95.1%, respectively.

1 Introduction As one of the earliest metals used by humans, copper is widely used in construction, defense, automobile manufacturing, power transmission and many other fields due to its excellent electrical and thermal conductivity [1–3]. Although pure copper has an excellent electrical and thermal conductivity, its hardness is low and the wear resistance is poor. It is difficult to meet the growing industrial demand, and its use range has been greatly limited. It has been found that the addition of a second phase reinforcement to the copper matrix can greatly improve the performance of the copper alloy [4], but the addition of most reinforcements tends to significantly reduce the conductivity of the composite [5]. Therefore, the selection of reasonable X. Ran (B) · Y. Wang · Y. Wang Key Laboratory of Advanced Structural Materials, Ministry of Education, Changchun University of Technology, Changchun 130012, Jilin, China e-mail: [email protected] Y. Wang e-mail: [email protected] Y. Wang e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_21

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reinforcements to develop copper-based composite materials with high strength and high conductivity has become an urgent problem for researchers. Since its discovery [6], graphene has been widely concerned by researchers for its excellent mechanical and physical properties. Its specific surface area is as high as 2600 m2 /g; its strength can reach 130 GPa, which is 100 times higher than that of steel [7]; its thermal conductivity up to 5000 W/m K [8] is the carbon material with the highest thermal conductivity so far, and its electron mobility can reach 15000 cm2 /V s at room temperature [9]. Studies have shown that the addition of graphene to the composite can effectively improve the overall performance of the composite. Hwang et. al [10] used a molecular-scale mixing method to prepare graphene copper-based composite materials by spark plasma sintering technology, which improved the interface between graphene and copper matrix, thereby improving the strength of composite materials; Kim et. al [11] used chemical vapor deposition to form a graphene film on the surface of copper substrate, and then electrochemically deposited a layer of copper on the graphene film. After repeated deposition, a graphene copper-based composite material with extremely high yield strength was obtained. Yang [12] used the mechanical stripping method to prepare graphene and used the spark plasma sintering method to obtain the graphene copper matrix composite with good mechanical properties. Graphene is an ideal two-dimensional nano-reinforcing material, but there are few reports on graphene/copper-based composites. Although the addition of graphene can improve the mechanical properties of the composite to a certain extent, it is often accompanied by a significant decrease in electrical conductivity and relative density [13–15]. This is mainly due to the poor bonding force between graphene and the matrix, which tends to agglomerate in the matrix and cannot be uniform. In order to improve the interfacial bonding ability, a high-strength- and high-conductivity graphene/copper-based composite material was prepared without sacrificing density and electrical conductivity. In this paper, a graphene/copper-based composite powder was prepared by in situ reduction method and discharged. The graphene/copper matrix composites were prepared by SPS process. The surface morphology and microstructure were characterized by SEM, XRD, and FR-IT. The effects of graphene content on the hardness, density and electrical conductivity of the composites were also studied.

2 Experimental Materials and Methods Materials Copper powder (>99.5%, 4–7 µm); flake graphite (>99.99%, 400 mesh); concentrated sulfuric acid, potassium permanganate, sodium nitrate, hydrogen peroxide, ascorbic acid, N-methyl pyrrolidone, hydrochloric acid, all of which are analytically pure. Preparation of Graphene Oxide In this experiment, a modified Hummers method was used to prepare graphene oxide [16]. The specific process was as follows: 1 g of

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graphite powder and 1 g of sodium nitrate were slowly added into a beaker containing 46 mL of concentrated sulfuric acid under stirring and mixed under ice-cooling for 15 min to mix it. Uniform; weighing 6 g of potassium permanganate slowly added to the mixed solution, stirring for 1 h in an ice bath to form a purple–green mixed solution; then, kept in a (35 ± 5) °C water bath pot for 1 h, while maintaining a moderate degree of stirring 40 mL deionized water was slowly added to the above beaker, and the reaction was vigorously observed. The reaction system rapidly increased in temperature, and a large amount of gas was released. The solution turned dark brown and the beaker was moved to a temperature of incubate in a (90 ± 5) °C water bath for 30 min. Then, add 100 mL of deionized water to dilute the solution to stop the reaction. Finally, add 6 mL of hydrogen peroxide with a mass fraction of 30%. Observe that the color of the solution turns bright yellow. After hot filtration, the filter cake was repeatedly washed with 3% hydrochloric acid and deionized water until the filtrate was detected with BaCl2 without a white precipitate, and then freeze-dried and collected for later use. Preparation of Graphene/Copper Composite Powders 0.2 g of graphene oxide was added to a beaker filled with 150 mL of a solution of N-methyl pyrrolidone to disperse ultrasonically for 1 h to uniformly disperse the graphene oxide in the solution; NaOH was added to the solution to adjust the pH to 9, and 2 g of ascorbic acid was added slowly. Add 25 mL of CuSO4 solution at a concentration of 0.1 mol/L dropwise. At this time, it can be observed that the color of the solution first turns to blue–brown and gradually turns black, and the beaker is moved to a water bath with a temperature of (90 ± 5) °C. Under heating and refluxing for 2 h, the product was filtered, washed with deionized water until the filtrate was detected with BaCl2 , and white precipitate was placed in a vacuum drying box for drying. Finally, graphene/copper composite powder was collected. Preparation of Graphene/Copper Composite Add 0.6, 1.2, and 1.8% volumetric graphene/copper composite powders to the copper powder and wet-mix the powders. Add the mixed powders to a graphite mold with a diameter of ϕ20 mm. The sintering parameters for the preparation of graphene/copper matrix composites are listed as follows: sintering temperature 700 °C, pressure 40 MPa, holding time 5 min, heating rate 100 °C/min. Finally a ϕ20, 3 mm thick round cake sample was obtained.

3 Results and Analysis Characterization of Graphene Oxide Figure 1 shows the characterization of graphene oxide. From Fig. 1a, we can see that the graphene oxide prepared by the modified Hummers method is dispersed in a thin yarn and partially transparent or translucent. The graphene has a smaller number of layers, and the XRD pattern in Fig. 1b shows that graphite has a high and sharp diffraction peak at 2θ of 26°. The query of the PDF card reveals that the diffraction peak is a single crystal graphite (002) plane. Compared with graphene oxide, there is no diffraction peak at this

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Fig. 1 Characterization of graphene oxide (GO) a TEM image b XRD pattern

angle but a wide and short diffraction peak appears around 2θ of 11°. According to the literature, the (002) crystal plane diffraction of graphite has been reported [17]. The peak shifts to the left as the degree of oxidation increases, and the diffraction peaks become shorter because the graphite oxide has an irregular crystal structure. Graphite’s spacing between layers of graphite under the influence of concentrated sulfuric acid and potassium permanganate strong oxidants. As a result, the crystal structure of the original rule was destroyed, crystallinity deteriorated, and disorder increased. The functional group possessed by graphene oxide is shown in Fig. 2. It can be seen from the figure that a large number of absorption peaks appear in the graphene oxide, and these absorption peaks correspond to the vibration of the hydroxyl group or the carboxyl group, respectively, which indicates that a large amount of inclusions are present in the graphene oxide structure. Oxygen functional groups, it is the presence of these functional groups that make graphene oxide more hydrophilic. When a reduction reaction occurs, Cu nanoparticles nucleate on the base and edges of graphene oxide due to the action of each epoxy group. The interfacial bonding ability of Cu and graphene is increased. Micromorphology of Graphene/Copper Composite Powders Figure 3a shows the SEM image of the graphene/copper composite powder. It can be seen from the figure that the reduced elemental copper particles are evenly distributed between the graphene sheets. This “infiltration” structure cannot only effectively increase the contact area between copper and graphene, thereby improving the interfacial bonding ability, and also effectively preventing the agglomeration of graphene. When the powder is mixed with copper powder, the element copper and copper powder distributed on the edge of the graphene sheet are effectively reduced. When fully mixed, graphene can be more uniformly dispersed in the copper matrix, and graphene and copper can be more fully combined. Figure 3b shows the XRD pattern of the composite powder. After comparison with the PDF card, the XRD image of the composite powder shows that the positions of the three strong peaks are coincident with the

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Fig. 2 FT-IR spectrum of graphene oxide (GO)

Fig. 3 Characterization of graphene/copper composite powders a SEM image b XRD pattern

face-centered cubic Cu, 2θ  43.3°, 50.4° and 74.1° diffraction. The peaks correspond to the (111), (200), and (220) crystal planes. The peak appearing around 23° in 2θ is the diffraction peak of graphene. Compare the XRD pattern of graphene oxide in Fig. 1b. The movement, consistent with the literature reports [17], proved that graphene oxide was reduced to graphene, and that in addition to the peaks of carbon and copper in the graph, no extra peak appeared, demonstrating the graphene prepared by the in situ reduction method. Among the copper powders, only graphene and elemental copper have no other impurity phases, and the reduction reaction is sufficient. Micromorphology of Graphene/Copper-Based Composite Figure 4 is the morphology and XRD patterns of graphene/copper-based composites with different graphene contents under optical electron microscope. It can be seen from the graph that the addition of graphene can significantly refine the grains because it is dispersed in copper. Graphene at the grain boundary of the matrix can act as a pinning effect, effectively prevent grain boundary migration, inhibit grain growth, and, with the increase of graphene content, the stronger the inhibition of grain growth, the smaller the size of the material grains. After the addition of graphene, the intensity and width

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Fig. 4 Morphology of graphene/copper-based composites with different graphene contents a 0 vol%, b 0.6 vol%, c 1.2 vol%, d 1.8 vol%, e XRD pattern

of the diffraction peaks of different graphene contents did not change significantly in the XRD pattern. This may be because the content of graphene in the composite powder is lower than the detection limit of XRD, so the characteristic peak to graphene is not been measured. Physical Properties of Composite Materials The hardness of the graphene/copper-based composite was tested using a microhardness tester. The

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Fig. 5 Effect of graphene content on microhardness

results are shown in Fig. 5. It was found that the overall microhardness of the composite material was higher than that of pure copper. When the content of graphene is 1.2%, the microhardness of the composite can reach 116 HV, which is 73% higher than that of pure copper (67 HV). According to the “fine grain strengthening” principle, this is because the addition of graphene can effectively suppress the growth of the copper matrix grains, refine the grains, and increase the strength of the composite materials. At the same time, the graphene also has a high strength, added to the copper matrix can play a role in the strengthening of the second phase. Graphene copper matrix composites prepared by the in situ reduction method can also make graphene more evenly dispersed between the graphene and copper matrix. The contact area is larger, the interface is stronger, and the reinforcement effect is better. Due to the non-interflow between graphene and copper, pores appear between the graphene and the copper matrix during the sintering process, which leads to a decrease in the relative density and electrical conductivity of the composite, and as the content of graphene increases, agglomeration occurs between graphene, which also leads to a decrease in the density and conductivity of the composite. Figure 6a, b show the variation of the relative density and conductivity of the composites with graphene content. It can be seen from the figure that the relative density and conductivity of the composites tend to decrease slightly with increasing graphene content. Consistent with the above description, when the content of graphene is 1.8%, the relative density of the composite material is as high as 98% or more, and the electrical conductivity is as high as about 95%. Compared with the previous research work (its IACS value is generally less than 80%), its relative density and electrical conductivity have reached very high values [18–20].

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Fig. 6 Effects of graphene content on relative density and electric conductivity

4 Conclusion (1) The graphene/copper composite powder prepared by the in situ reduction method can uniformly disperse the elemental copper between the graphene sheets, effectively inhibit agglomeration between the graphene sheets and enhance the interface binding ability between the graphene and the matrix. (2) The graphene/copper-based composite material has been successfully prepared by spark plasma sintering, and graphene can be uniformly dispersed in the copper matrix without obvious agglomeration. (3) The hardness of graphene/copper-based composites is significantly higher than that of pure copper. When the volume fraction of graphene is 1.2%, the hardness of the composite is 116 HV, which is 73% higher than that of pure copper. The density and conductivity of the composites decrease slightly with the increase of graphene, but the decrease is small. The density and conductivity of the composites are 98 and 95%, respectively, demonstrating that the interface between the composite powder graphene and copper base is well-bonded.

References 1. H.J. Zhao, L. Liu, Y.T. Wu, W.B. Hu, Investigation on wear and corrosion behavior of Cugraphite composites prepared by electroforming. Compos. Sci. Technol. 67, 1210–1217 (2007) 2. R.R. Zahran, I.H.M. Ibrahim, G.H. Sedahmed, The corrosion of graphite/copper composites in different aqueous environments. Mater. Lett. 28, 237–244 (1996) 3. P.K. Rohatgi, S. Ray, Y. Liu, Tribological properties of metal matrix-graphite particle composites. Int. Mater. Rev. 37(3), 129–149 (1992) 4. D.L. Wang, Y. Feng, S. Li et al., Preparation and properties of alumina dispersion strengthened copper matrix composites. Met. Funct. Mater. 16(2), 21–25 (2009) 5. P. Zhang, J. Jie, H. Li et al., Microstructure and properties of TiB2 particles reinforced Cu-Cr matrix composite. J. Mater. Sci. 50(9), 3320–3328 (2015)

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6. K.S. Novoselov, A.K. Geim, S.V. Morozov et al., Electric field effect in atomically thin carbon films. Science 306(5296), 666 (2004) 7. J.H. Chen, C. Jang, S.D. Xiao et al., Intrinsic and extrinsic performance limits of graphene devices on SiO2 . Nat. Nanotechnol. 3(4), 206–209 (2008) 8. A.A. Balandin, S. Chosh, W.Z. Bao et al., Superior thermal conductivity of single-layer grapheme. Nano Lett. 8(3), 902–907 (2008) 9. Service R F, Carbon sheets an atom thick give rise to graphene dream. Science 324(59291), 875–877 (2009) 10. J. Hwang, T. Yoon, S.H. Jin et al., Enhanced mechanical properties of graphene/copper nanocomposites using a molecular-level mixing process. Adv. Mater. 25(46), 6724–6729 (2013) 11. Y. Kim, J. Lee, M.S. Yeom et al., Strengthening effect of single-atomic-layer graphene in metal-graphene nanolayered composites. Nat. Commun. 4, 2114 (2013) 12. S. Yang, Preparation and Properties of Less Layer Graphene Reinforced Copper Matrix Composites (Harbin Institute of Technology, 2011) 13. J. Kováˇcik, Š. Emmer, J. Bielek, Effect of composition on friction coefficient of Cu–graphite composites. Wear 265(3), 417–421 (2008) 14. T. Futami, M. Ohira, H. Muto, M. Sakai, Contact/scratch-induced surface deformation and damage of copper–graphite particulate composites. Carbon 47(11), 2742–2751 (2009) 15. F. Akhlaghi, A. Zare-Bidaki, Influence of graphite content on the dry sliding and oil impregnated sliding wear behavior of Al 2024–graphite composites produced by in situ powder metallurgy method. Wear 266(1), 37–45 (2009) 16. L.J. Cote, F. Kim, J. Huang, Langmuir-Blodgett assembly of graphite oxide single layers. Am. Chem. Soc. J. 131(3) (2008) 17. K.-C. Tsai, H.-C. Kuan, H.-W. Chou et al., Preparation of expandable graphite using a hydrothermal method and flame-retardant properties of its halogen-flee flame-retardant HDPE composites. Polym. Res. 18, 483–488 (2011) 18. A.D. Moghadam, E. Omrani, P.L. Menezes et al., Mechanical and tribological properties of self-lubricating metal matrix nanocomposites reinforced by carbon nanotubes (CNTs) and graphene—a review. Compos. B 77, 402 (2015) 19. T. Varol, A. Canakci, Microstructure electrical conductivity and hardness of multi-layer graphene/copper nanocomposites syn the sized by flake powder metallurgy. Met. Mater. Int. 21(4), 704 (2015) 20. G. Huang, H. Wang, P. Cheng et al., Preparation and characterization of the graphene-Cu composite film by electro deposition process. Microelectron. Eng. 157, 7 (2016)

The Effect of Aluminum Content on the Microstructures of Single-Phase γ-TiAl-Based Alloy Yaodong Xuanyuan, Yan Long, Yinbiao Yan and Sen Yang

Abstract TiAl-based alloys have not only excellent high-temperature strength, hightemperature creep and oxidation resistance, but also low density and high strength-toweight ratio, having extensive application prospect in the aerospace and automobile industry. In this paper, the single-phase γ-TiAl-based alloys were studied. In order to refine the coarse microstructures of single-phase γ-TiAl-based alloys, three different kinds of γ-TiAl-based alloys, including Ti–54Al (at.%), Ti–52Al (at.%), and Ti–50Al (at.%), were melted and homogenized at 1150 °C for 36 h. The effect of composition on grain refinement and grain boundary distribution of single-phase γ-TiAl-based alloys were investigated. The results showed that the grain size of the alloys was refined with the increase of aluminum content after homogenizing treatment. In addition, the grain boundary character distribution was non-uniform. Inside the columnar crystals, there were a large number of coherent 3 grain boundaries. The derived capacity of 3 grain boundaries was poor, and general high-angle grain boundaries were connected to each other.

1 Introduction γ-TiAl has attracted a great deal of attention in the field of aerospace, and also recently in the automobile industry, because of its potentially attractive properties for high-temperature structural applications such as low density, good high-temperature strength, and superior resistance to oxidation (above 750 °C) [1, 2]. In the past two decades, the composition and phase relationship of TiAl alloys are roughly determined, as shown in Fig. 1 [3]. Among them, γ-TiAl-based alloys are a hot area of research in TiAl alloys. γ-TiAl-based alloys include single-phase alloys and duplex alloys. Alloys containing 52–56% aluminum are γ-TiAl single-phase structures. The γ phase is the structure of L10, which is close to the face-centered cubic ordered Y. Xuanyuan · Y. Long · Y. Yan · S. Yang (B) School of Material Science and Engineering, Nanjing University of Science and Technology, No. 200 Xiaolingwei Street, Nanjing 210094, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_22

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Fig. 1 Ti–Al binary phase diagram [3]

structure. Ti atoms and Al atoms are alternately arranged on the (002) crystal plane, and the axis ratio of a/c is generally not 1. In the past few decades, Watanabe introduced the concept of “an approach to grain boundary engineering design and control for strong and ductile polycrystals” in 1984 [4], which has been developed as grain boundary engineering (GBE) by Palumbo [5–9]. The performance enhancement associated with GBE is attributed to the increase in the proportion of low- (1 ≤  ≤ 29) coincidence site lattice (CSL) boundaries and redistribution of grain boundary network topology. In recent years, the grain boundary engineering of annealing twins is mainly applied to the middlelow stacking fault energy of the structure of FCC, which are austenitic stainless steel alloy and nickel-based alloy. The γ phase with the structure of L10 in γ-TiAl-based alloys is very close to FCC, also has low stacking fault energy. Theoretically, it is suitable for grain boundary engineering of annealing twins.

2 Materials and Experiment According to the composition of the experimental design, the experiment has three kinds of γ-TiAl-based alloy, including Ti–54Al (at.%), Ti–52Al (at.%), and Ti–50Al (at.%), which depended on the aluminum content. The alloy was melted into ingots in an induction melting furnace, and each ingot was repeatedly melted four times. Alloys of each component were melted two to five ingots with a mass of about 50 g. Afterward, three kinds of the γ-TiAl-based alloy were homogenized at 1150 °C for 36 h. The optical microstructure observations were conducted by Carl Zeiss AxioVision microscope. Hitachi SU1510 Tungsten Filament Gun Scanning Electron

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Microscope (SEM) equipped with an Oxford Instruments Nordlys EBSD system was used to quantify the average intragranular misorientations and determine the character of grain boundaries.

3 Result 3.1 The Effect of Aluminum Content on the Microstructures of As-cast Alloys According to the Ti–Al binary phase diagram, when the aluminum content is 50–56 at.%, the alloy is single-phase γ-TiAl microstructure. In order to obtain the single-phase γ-TiAl alloys, three kinds of alloys were designed: Ti–54Al, Ti–52Al and Ti–50Al. The microstructures of the three different components of as-cast alloys are shown in Fig. 2. All microstructures showed obvious dendrite morphology. In the figure, the black area is dendritic morphology, and the white area is the aluminumrich γ segregation area. As the aluminum content decreased, the dendrite volume increased and the γ segregation area gradually decreased. A large number of columnar crystals could be seen in the Ti–54Al alloys, and twins existed along the long axis. These columnar crystals were formed when the alloys were smelted and directly cooled by the γ liquid phase. However, no similar columnar crystals existed in Ti–52Al and Ti–50Al. The XRD results of three different kinds of as-cast alloys are shown in Fig. 3. XRD pattern reveals that there was mainly γ phase in the Ti–54Al and Ti–52Al alloys, only a small amount of α2 phase existed in the Ti–54Al alloys. The composition of the Ti–50Al alloy is in the γ single-phase region of the phase diagram. Nevertheless, a certain amount of α2 -Ti3Al phase still existed when the aluminum content was 50 at.% because of its being at the edge of the phase zone and the segregation of as-cast alloy.

3.2 The Effect of Aluminum Content on Microstructures and Grain Boundary Character Distribution of Homogenized Annealed Alloys The as-cast microstructure of γ-TiAl alloy was not uniform and there was obvious segregation. It should be homogenized before actual machining deformation. Shorttime homogenization is difficult to eliminate dendrites, so long-time treatment with high temperature was used. Figure 4 shows that the three-component samples were homogenized at 1150 °C for 36 h. It can be seen from Fig. 4a that the annealed microstructures of Ti–54Al are mainly columnar crystals and a small number of equiaxed grains. The average length of the long axis columnar crystals was about 600 μm, and twin crystals were formed along the long axis direction, and twins

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Fig. 2 Microstructures of γ-TiAl-based as-cast alloy a Ti–54Al; b Ti–52Al; c Ti–50Al

Fig. 3 X-ray diffraction pattern of γ-TiAl-based alloy a Ti–54Al; b Ti–52Al; c Ti–50Al

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were formed along the long axis direction, while the equiaxed grains were mainly formed by annealing, and the microstructure was extremely non-uniform. Equiaxed γ single-phase microstructures were formed by annealing in the Ti–52Al alloys. The annealed microstructure of Ti–50Al was also in the form of single-phase γ-TiAl alloy. No obvious lamellar grains existed, and the average grain size was about 250 μm. However, some researches [10–12] have shown that in the range of 50–52 at.% of aluminum content, the alloy may still contain a small amount of α2 phase after annealing, and it is not completely a single-phase γ-TiAl alloy. By comparing the microstructures of three kinds of annealed alloys with different aluminum contents, it could be found that with the increasing of aluminum content, the grain size was increased obviously, and the single-phase γ microstructures represented by Ti–54Al had the largest grain size and contained a large amount of coarse columnar grains. At the same time, there were also equiaxed grains formed by annealing, and the microstructures were not extremely uniform. The room temperature properties of γ-TiAl alloys are considered to be strongly dependent on their microstructures. The microstructures, which were coarse and nonhomogeneous, tended to affect the overall mechanical properties of the alloys. Refinement of grains is the best way to improve their overall performance. According to the Hall–Petch relationship, refining the grains can effectively improve the slip resistance of the material. Meanwhile, with the increase of grain boundaries, the stress generated during plastic deformation could be sustained by more grains, which can also enhance the coordination of the grains. Table 1 shows the grain boundary character distribution of the annealed alloys with different compositions. The influence of the composition on the grain boundary character distribution of γ-TiAl alloys can be seen. The special grain boundaries of the three kinds of the alloys were dominated by 3 grain boundaries, the Ti–50Al alloys had a higher ratio of specific grain boundaries. With the increase of aluminum content, the specific grain boundary ratio of Ti–52Al alloys decreased, while the specific grain boundary ration of Ti–54Al alloys was higher than that of Ti–52Al alloys, the main reason was that had a higher ratio of 3 grain boundaries. In addition, the ratio of f (9 + 27)/f (3) is used to describe the ability of the unit length of the 3 grain boundaries to derive the 9, 27 grain boundaries. The ability of 3 grain boundaries to derive the of 9 and 27 grain boundaries by migration was stronger with the increase of the value, forming a large number of trigeminal grain boundaries composed of 3n grain boundaries [13, 14]. It can be seen from the table that the value of f (9 + 27)/f (3) of the Ti–54Al alloy is extremely low, indicating that the ability of derivatization is very low.

Table 1 The fraction of low-CSL boundaries in three kinds of γ-TiAl alloys (%) Sample

3

9

27

3n

Total-

f (9 + 27)/f (3)

Ti–54Al

33.18

0.25

0.17

33.60

35.92

1.27

Ti–52Al

29.01

0.99

0.39

30.39

34.13

4.76

Ti–50Al

39.08

1.44

0.27

40.79

44.19

4.38

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Fig. 4 Microstructures of γ-TiAl-based alloy after annealing at 1150 °C for 36 h a Ti–54Al; b Ti–52Al; c Ti–50Al

Figure 5 shows the orientation maps and EBSD maps of three kinds of microstructures. The red, blue, green, and black lines in Fig. 5 represent the 3 boundary, 9 boundary, 27 boundary and random boundary respectively. In the study of grain boundary character optimization, it is not sufficient to determine whether the grain boundary has been optimized only by increasing the proportion of the special grain boundaries. It is also necessary to consider the breaking effect of the special grain boundary on the connectivity of general large-angle grain boundaries. Because of the general large-angle boundaries connected with each other, even if the ratio of the special boundaries increases, the failure crack will continue to expand along the general large-angle boundaries. There were a large number of columnar crystals and a small number of equiaxed grains in the scanning region of Ti–54Al alloys. It was obvious that the columnar crystals contained more coherent 3 grain boundaries, but mainly distributed along the long axis direction of the columnar crystals. The grains were completely surrounded by generally large-angle grain boundaries, and the special grain boundaries were less in the equiaxed grains. Figure 5c, f, i is grain boundary reconstruction maps. It can be observed that the special grain boundaries in the Ti–54Al alloy had almost no breaking effect on the general high-angle grain boundaries, and there was no optimization on the grain boundaries. With the aluminum content decreasing, the connectivity of the general high-angle grain boundaries was interrupted.

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Fig. 5 Orientation maps and EBSD maps of annealed γ-TiAl alloys a, b, c Ti–54Al; d, e, f Ti–52Al; g, h, i Ti–50Al

By analyzing the ratio of special grain boundaries of Ti–54Al alloy, it is found that the columnar crystals contained a higher proportion of special grain boundaries. In order to accurately analyze the grain boundary character distribution of the columnar region of Ti–54Al alloy, a smaller area is selected, as shown in Fig. 6. A columnar often contained several coherent 3 grain boundaries, as shown in Fig. 6b, which makes the Ti–54Al alloys exhibited a higher ratio of specific grain boundaries, but had almost no interruption to the general high-angle grain boundaries. As a result, these growth twins were completely surrounded by general high-angle grain boundaries, as shown in Fig. 6c. Moreover, the ratio of specific grain boundaries of the columnar regions after accurate calibration reaches 38.48%. Due to the limitation of EBSD accuracy, a small amount of extremely fine twins cannot be calibrated, but the grain boundary engineering mainly studies the effect

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Fig. 6 Orientation maps (a) and grain boundary reconstruction maps (b, c) of the columnar region of Ti–54Al alloys

of annealing twins. These extremely fine growth twins along the long axis did not optimize the grain boundary character distribution.

4 Conclusion In this paper, the effect of aluminum content on grain refinement and grain boundary character distribution of single-phase γ-TiAl alloys has been studied. After homogenizing treatment, the grain size of the alloys is greatly improved with the increase of aluminum content. In particular, the microstructures of Ti–54Al consist of coarse columnar crystals and a small number of equiaxed grains, which show inhomogeneously. The grain boundary character distribution is also non-uniform, and the columnar crystals contain a large number of coherent 3 grain boundaries. The derived capacity of 3 grain boundaries is poor, and general high-angle grain boundaries are connected to each other.

References 1. G. Welsch, R. Boyer, E.W. Collings, Materials Properties Handbook: Titanium Alloys (ASM international, 1993) 2. F. Appel, J.D.H. Paul, M. Oehring, Gamma Titanium Aluminide Alloys: Science and Technology (Wiley, 2011) 3. J.C. Schuster, M. Palm, Reassessment of the binary aluminum-titanium phase diagram. J. Phase Equilib. Diffus. 27(3), 255–277 (2006) 4. T. Watanabe, An approach to grain-boundary design for strong and ductile polycrystals. Res. Mech. 11(1), 47–84 (1984) 5. G. Palumbo, E.M. Lehockey, P. Lin, Applications for grain boundary engineered materials. JOM 50(2), 40–43 (1998) 6. C. Cheung, U. Erb, G. Palumbo, Application of grain boundary engineering concepts to alleviate intergranular cracking in alloys 600 and 690. Mater. Sci. Eng. A 185(1–2), 39–43 (1994)

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7. E.M. Lehockey, G. Palumbo, P. Lin, Improving the weldability and service performance of nickel-and iron-based superalloys by grain boundary engineering. Metall. Mater. Trans. A 29(12), 3069–3079 (1998) 8. G. Palumbo, Thermomechanical Processing of Metallic Materials. U.S. Patent No. 5,702,543 (1997) 9. G. Palumbo, Metal Alloys Having Improved Resistance to Intergranular Stress Corrosion Cracking. U.S. Patent No. 5,817,193 (1998) 10. D. Shechtman, M.J. Blackburn, H.A. Lipsitt, The plastic deformation of TiAl. Metall. Trans. 5(6), 1373–1381 (1974) 11. R.G. Baligidad, U. Prakash, V.R. Rao et al., Effect of addition of β-phase stabilizing elements (Nb, Mo and V) on plastic behaviour of Ti3 Al single crystals with the D019 structure. Metall. Trans. A 6(11), 1991 (1975) 12. G. Hug, A. Loiseau, A. Lasalmonie, Nature and dissociation of the dislocations in TiAl deformed at room temperature. Philos. Mag. A 54(1), 47–65 (1986) 13. V. Randle, Mechanism of twinning-induced grain boundary engineering in low stacking-fault energy materials. Acta Mater. 47(15), 4187–4196 (1999) 14. V. Randle, Twinning-related grain boundary engineering. Acta Mater. 52(14), 4067–4081 (2004)

Effect of Al–Er–Zr Master Alloy on Grain Refinement After Heat Treatment Haiyue Yu, Hui Huang, Zuoren Nie, Shengping Wen, Kunyuan Gao and Wei Wang

Abstract The effects of Al–0.5Er–0.2Zr and Al–1Er–0.25Zr (wt%) alloys on the grain size of high-purity aluminum after isothermal aging at 450 °C were studied by micro-hardness, scanning electron microscopy (SEM), optical microscopy, and transmission electron microscope (TEM). The experimental results showed that the Al3 Er secondary phases precipitated uniformly after isothermal aging of Al–0.5Er–0.2Zr alloy. The content of Er element in Al–1Er–0.25Zr alloy was much larger than the maximum equilibrium solid solubility in aluminum. The primary phase was densely distributed in the matrix. In addition, a large number of Al3 Er and Al3 (Er, Zr) phases matched with the matrix lattice have been precipitated during heat treatment. The heat-treated master alloys were added to the high-purity aluminum liquid, and the nucleation sites were provided by the micron-sized and nano-sized precipitation phases. At last, the average grain size obtained by the refinement experiment was less than 300 µm. At the same time, the reasons for the refinement effect of the interaction between Er and Zr elements were preliminarily analyzed.

H. Yu (B) · H. Huang · Z. Nie · S. Wen · K. Gao · W. Wang School of Materials Science and Engineering, Beijing University of Technology, Beijing 100124, People’s Republic of China e-mail: [email protected] H. Huang e-mail: [email protected] Z. Nie e-mail: [email protected] S. Wen e-mail: [email protected] K. Gao e-mail: [email protected] W. Wang e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_23

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1 Introduction Aluminum alloys have been widely used in various fields because of their many advantages in performance. However, aluminum and its alloys tend to form coarse grain structures under the casting conditions, which have a great influence on the properties. Currently, it is commonly used to refine the grain at casting stage. Grain refinement is of great significance to the semi-continuous casting of deformed aluminum alloy and in the casting of cast aluminum alloys. The equiaxed grain structure is beneficial to the improvement of mechanical properties. It can reduce segregation, reduce the tendency of thermal cracking, improve the shrinkage, elimination or better dispersion of the castings in the solidification process, improve the airtightness and the surface of the castings [1, 2]. Al–Ti–B master alloy has been developed for nearly 60 years. About 75% of the world aluminum industry uses Al–Ti–B to refine grain [3]. The refiner contains soluble TiAl3 and insoluble TiB2 particles to increase the nucleation position and achieve nucleation in the melt. However, there are still some problems in Al–Ti–B, such as boride agglomeration, low efficiency of refinement [4], and some elements such as Zr, V, Cr poisoning [5, 6]. So far, there is not a kind of viewpoint and theory that can completely explain all the grain refinement phenomena. With the in-depth understanding of the refining mechanism, the effective grain refinement needs the good lattice matching of the solid particles and the matrix. The refiner particles are evenly distributed in the liquid phase, and there must be sufficient particle number density. These views are widely accepted as the theoretical basis for the research and development of the new refiner [7]. An in-depth study of rare earth Er elements by our research group found that the addition of Er element in aluminum alloy can refine the grain size of the as-cast alloy, improve the strength and the comprehensive performance of the alloy [8, 9]. Moreover, nano-sized Al3 Er particles have a good coherent relationship with the matrix [10, 11]. The addition of Zr can not only significantly improve the microstructure of aluminum alloy, but also refine the alloy grain [12]. The interaction of Er and Zr will increase the stability of the whole alloy [13–15]. In this paper, Al–Er–Zr is used as master alloy to be added to high-purity aluminum after heat treatment. The effects of precipitates on grain refinement are explored.

2 Materials and Methods In this experiment, the Al–Er–Zr master alloy was prepared by the traditional ingot metallurgy method. The experimental alloys were prepared by high pure aluminum (99.99 wt%) with Al–6Er and Al–4Zr (wt%) master alloys, and the alloy composition ratio is shown in Table 1. In a resistance crucible furnace, a graphite clay crucible was used to melt the high-purity aluminum at 780–800 °C, then Al–Er and Al–Zr master alloys were added, and the temperature was heated to 900 °C for

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Samples

Er

Zr

Al

1

0.5

0.20

Balance

2

1.0

0.25

Balance

2 h. This aim was to decompose the primary phase in the master alloy. After stirred evenly, the liquid metal was poured into a cold iron mold of 5 × 65 × 180 mm3 , and the pouring temperature was 700 °C (the purpose of making the thin ingot is to obtain a higher cooling rate for the ingot, increase the solid solubility of the element, and enhance the heat treatment precipitation effect). After the ingot had solidified, it was quickly quenched to room temperature. The Al–Er–Zr master alloy were aged at 150–600 °C with increments of 50 °C, each temperature lasting 2 h. After that, the isothermal aging temperature was 450 °C, and the aging time was 24 h. After each aging step, precipitation was measured by Vickers hardness. Each sample used an average of at least 10 independent measurements with a load of 100 g and a dwell time of 10 s. The experiment used a JEOLJSM-6500F scanning microscope from JEOL Ltd., and the working voltage of the scanning electron microscope was set to 20 kV. Transmission electron microscopy (TEM) sample sections were ground to less than 70 µm and punched into 3 mm discs. The thin foils were then prepared for TEM observation by twin-jet polishing with an electrolyte solution consisting of 30% HNO3 and 70% methanol below −30 °C. The specimens were analyzed using a JEOL 2010 microscope with an operating voltage of 200 kV. The method of adding the Al–Er–Zr master alloy is as follows: at first, the highpurity aluminum was melted at 800 °C and kept for one hour. When the aluminum liquid was cooled down to 700 °C, the Al–Er–Zr master alloy was added. The amount added is 0.1% of the matrix according to the Er element, and the temperature was kept for 3 min, fully stirred and poured. The refined samples were corroded with Keller reagent (0.5%HF + 1.5%HCl + 2.5%HNO3 + 95.5%H2 O) to observe the grain size of the alloy. Because the refined ingot is a cylindrical ingot with a diameter of 60 mm and a height of 25 mm, the grain size statistical area was selected to be fixed at a position 5 mm from the center.

3 Results and Discussion Figure 1a shows the relationship between temperature and Vickers hardness in isochronous aging process of Al–0.5Er–0.2Zr and Al–1Er–0.25Zr alloys. Due to the content of alloy components, the hardness of the Al–1Er–0.25Zr alloy is almost slightly higher than that of the Al–0.5Er–0.2Zr alloy. The first peak hardness of Al–0.5Er–0.2Zr occurred at 350–400 °C. The first peak hardness of Al–1Er–0.25Zr alloy appeared at 250–350 °C. Both alloys reach the peak hardness point at 500–550 °C. The hardness of Al–0.5Er–0.2Zr alloy is 34.3 HV; the hardness value of

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Fig. 1 a Al–Er–Zr master alloy isochronous aging, temperature-Vickers hardness diagram; b Al–Er–Zr master alloy 450 °C isothermal aging 24 h, time-Vickers hardness diagram

Al–1Er–0.25Zr is 35.0 HV, then hardness values decreases. In addition, the hardness of Al–0.5Er–0.2Zr alloy with lower composition content is the most obvious, which is about 9 HV higher than that in the as-cast state. The hardness value decreases sharply between 550 and 600 °C. In the isochronous process, the peak hardness at 550 °C should be lower than this temperature when the isothermal aging temperature is selected. This paper selects the isothermal aging at 450 °C. Figure 1b is the relationship between time and Vickers hardness of Al–0.5Er–0.2Zr and Al–1Er–0.25Zr alloys at 450 °C isothermal aging for 24 h. It can be seen that the hardness peaks of the two components in the alloy within 24 h of aging. It is worth noting that the Al–1Er–0.25Zr alloy with higher component content has two hardness peaks. The hardness of Al–0.5Er–0.2Zr alloy increases at the initial stage of aging, and then reaches a peak hardness of 41.3 HV at 5 h. The precipitated phase at this time is evenly and densely distributed. At the same time, there is a great effect of increasing the relative hardness of nanoscale precipitation. As time goes by, the precipitated phase gradually coarsens, which leads to the decrease of hardness. The Al–1Er–0.25Zr alloy shows a small hardness peak at 1 h after aging. It was found that the hardness of the alloy was strengthened because of the large precipitation of the Al3 Er phase, and the subsequent decrease in hardness is due to the coarsening of Al3 Er phase [15]. When the aging time reaches 15 h, a second peak hardness of 45.6 HV occurs. This is because the Zr element slowly accumulates on the nanometer Al3 Er phase to form a large number of nanometer Al3 (Er, Zr) composite phases. The phase of this L12 structure is a typical coreshell structure, which can enhance the matrix. Although the phase of the core-shell structure is very stable, due to the existence of a large number of primary phases in the alloy matrix, this strengthening effect is gradually masked by the coarsening of the primary phase over time, which reduces the hardness value of the alloy. The aging strengthening is manifested by the analysis of the second phase. Figure 2 is a scanned image of the alloy’s as-cast and isothermal aging hardness peaks. The

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Fig. 2 SEM of different states of Al–Er–Zr master alloy. a Al–0.5Er–0.2Zr alloy as-cast microstructure; b Al–0.5Er–0.2Zr alloy isothermal aging for 5 h microstructure (peak point of hardness); c Al–1Er–0.25Zr alloy as-cast microstructure; d Al–1Er–0.25Zr alloy isothermal aging for 15 h microstructure (peak point of hardness)

design of the two alloy’s elements exceeds the maximum equilibrium solid solubility, so there are a large number of primary phases in the as-cast state, and these primary phases contain the Er element, indicating that the supersaturated solid solubility of Er in the aluminum matrix is very limited. No primary phase containing Zr is found in the alloy. Because of the rapid cooling rate under the experimental conditions, the Zr element is completely dissolved in the aluminum matrix to form a supersaturated single-phase solid solution, which is beneficial to the precipitation of the aging strengthening phase. It is observed from Fig. 2a that the precipitation phase distribution in the cast state of Al–0.5Er–0.2Zr alloy is uneven, and there is a large number of spherical particles, while some region is mainly precipitated along the grain boundary. The peak point of hardness reached at 5 h after isothermal aging. The precipitated phase is Al3 Er evenly distributed and spherical particles are precipitated densely. The segregation phenomenon is completely eliminated (Fig. 2b). In Fig. 2c, the primary phase of Al–1Er–0.25Zr alloy is distributed in the crystal and grain boundary, and most of the primary phase precipitates at the grain boundary, showing a bone shaped intermittent or continuous distribution. After 15 h of isothermal aging, a large number of spherical Al3 Er phases precipitated in the grain, and the primary phase of the grain boundaries is spheroidized and decreased (as shown in Fig. 2d).

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Fig. 3 TEM of Al–Er–Zr master alloy at 450 °C isothermal aging hardness peak state. a Al–0.5Er–0.2Zr alloy isothermal aging 5 h (hardness peak point) transmission diagram; b Al–1Er–0.25Zr alloy isothermal aging 15 h (hardness peak point) transmission diagram

In addition, TEM observation of the peak point of isothermal aging at 450 °C shows that there are a large number of nanoscale precipitates (Fig. 3). Figure 3a is a transmission picture of isothermal aging of Al–0.5Er–0.2Zr alloy for 5 h, and bean petal precipitates can be observed. The size is between 10 and 15 nm and the previous study shows that the particle is the Al3 Er phase of the L12 structure formed in the supersaturated solid solution, and it maintains a coherent relationship with the matrix [16–18]. These small Al3 Er particles are evenly distributed in the matrix and have a good pinning effect on the dislocation, so the alloy has obvious aging strengthening effect. The transmission picture of Al–1Er–0.25Zr alloy is shown in Fig. 3b. The precipitated phase has a particle size of 15–20 nm, a large amount and evenly dispersion distribution, and is coherent with the aluminum matrix. By transmission electron microscopy and energy spectrum analysis, the nanoparticles contain 3 elements of Al, Er and Zr in addition to Al3 Er, and the particles are Al3 (Er1-x Zrx ) composite phase combined with the literature reported in the literature [19]. The L12 -type structure is effective in reinforcing the dislocation and subgrain boundaries. It is an effective strengthening phase in aluminum alloys. Because the different atomic radius of the Er and Zr elements, the nuances of the lattice structure in different regions inside the ternary particles lead to the difference in diffraction contrast (lattice distortion). The diffusion rate of Er atom is fast and lies in the core part of the phase. The diffusion rate of Zr atom is slow and attached to the periphery of the Er phase. The phase of the core-shell structure is stable at high temperature, which is the reason for the strengthening of the aging. On the basis of the above experiments, refinement experiments use the as-cast state of Al–0.5Er–0.2Zr master alloy, the isothermal aging treatment at 450 °C for 5 h and the 10-h overaging state. The Al–1Er–0.25Zr master alloy use the as-cast, isothermal aging treatment at 450 °C for 10 h and hardness peak at 15 h. The peak point of hardness in isothermal aging is selected because it is relatively small in size and more evenly distributed. At the same time, refinement experiments are performed on master alloys that are not in the peak state of aging in order to make a comparison.

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Figure 4a is the refinement result of Al–0.5Er–0.2Zr master alloy at the casting state, and the result of the refinement is cooling by water. Because the cooling rate is an important factor affecting grain size, the size of the grain is very small after refinement. The average grain size is 125 µm, and it has a large number of equiaxed grains. Figure 4c shows the refinement of the Al–1Er–0.25Zr master alloy in the as-cast state. This result also uses water cooling to reduce the melt temperature, but the grain size is very large at 275 µm. This is because the precipitated phases in the as-cast state are distributed along the grain boundaries and are not easily regarded as effective nucleation sites. The reason for refinement is the segregation of Er at grain boundaries to form the primary Al3 Er phase, which hinders grain growth. The grain refinement produced by the concentration gradient mechanism. Although the cooling rate has a great influence on the grain size in the casting process, the heterostructure nucleus needs an effective nucleation site, and the Al–0.5Er–0.2Zr master alloy has a large number of spherical Al3 Er phases in the as-cast state, which is similar to the TiB2 particles in the Al–Ti–B refiner, that is easy to promote nucleation. This is why the refining effect of the Al–1Er–0.25Zr master alloy is worse than the Al–0.5Er–0.2Zr master alloy. But b and d in Fig. 4 are air-cooled. The purpose is to show the effect of precipitated particles on grain refinement. Figure 5 shows the average size of the micron-sized precipitates observed under scanning electron microscopy and the corresponding average size of the refined grains after acting as a refiner. It can make a full discussion on the refining effect of Al–Er–Zr master alloy. For the Al–0.5Er–0.2Zr master alloy, the size of the precipitated phase has a certain effect on the grain refinement. The smaller the particle size, the smaller the grain size obtained. The average particle size of the primary phase in the as-cast condition of Al–0.5Er–0.2Zr master alloy is 1.82 µm (Fig. 5a, 1). With the aging process, the precipitated phase is coarsening. The peak state of hardness is reached at 5 h with isothermal aging, the average particle size is 1.95 µm (Fig. 5a, 2), and the corresponding grain size is 180 µm (Fig. 5b, 2). The average particle size is 2.08 µm for 10 h of aging, and the corresponding grain refinement results is 240 µm (Fig. 5b, 3). Because the Al–0.5Er–0.2Zr master alloy has a large number of spherical particles in the as-cast and aging conditions, and the precipitated nano-sized Al3 Er phases after heat treatment are lattice-matched with the matrix, they are easy to act as nucleating agents. Eliminates the condition of large-sized grains that have appeared due to the fact that there is no temperature reduces in water-cooled conditions. The cooling rate and the utilization of nucleating agent in the refining process have a great influence on the grain size, and these factors cannot be ignored. For Al–1Er–0.25Zr master alloy, 10 h of isothermal aging does not reach its peak hardness. Most of the precipitates still accumulate in the grain boundary, and there are few spherical particles in the micron grade. But by the heat treatment results of Al–0.5Er–0.2Zr master alloy, it is known that the nano-phase of Al–1Er–0.25Zr master alloy must exist, and the larger concentration gradient during the solidification process limits the grain growth. The average particle size is 255 µm (Fig. 5b, 5). After isothermal aging for 15 h, a large amount of micron-secondary phase Al3 Er are precipitated, and the grain boundary precipitates are melted and dispersed to obtain an average particle size of 2.0 µm (Fig. 5a, 6). At the same time, There are a large

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Fig. 4 Metallographic picture of Al–Er–Zr master alloy added to high-purity aluminum. a Al–0.5Er–0.2Zr alloy is added to the high-purity aluminum refining result under the as-cast condition (cooling in water); b Al–0.5Er–0.2Zr alloy is added to the high-purity aluminum refining result under isothermal aging for 5 h; c Al–1Er–0.25Zr alloy is added with high-purity aluminum refining result under as-cast conditions (cooling in water); d Al–1Er–0.25Zr alloy is added to highpurity aluminum refining result under isothermal aging for 15 h. (Scaler 500 µm)

Fig. 5 a Statistics of precipitate phases in the micron level for different states of the Al–Er–Zr master alloy; b grain size after refinement obtained from the master alloy in different states

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number of 15–20 nm Al3 (Er1-x Zrx ) composite phase. The composite phase together with the micron-sized Al3 Er phase acts as a refinement. The average grain size is 179 µm (Fig. 5b, 6). In addition, the size of the precipitated phase in the micron level of the master alloys of different compositions is not comparable with the obtained grain size, because the higher the composition, the greater effect of grain boundary precipitation phase hinders the grain growth. The nano-precipitate phase after heat treatment will also affect the grain size. So the influence factors of grain size are not the single particle size.

4 Conclusions Through the above experimental analysis, we can draw the following conclusions: 1. Al–Er–Zr master alloy can be used as an effective grain refiner. The primary phase Al3 Er presents in the master alloy, and the secondary phase Al3 Er, Al3 (Er, Zr) composite phase precipitate during the aging process. Because the large number of particles are evenly distributed and have good coherency with the aluminum matrix, it can be used as an effective nucleation particle to refine the grain. 2. The size and morphology of precipitated phase in the same composition of Al–Er–Zr master alloy have a certain influence on the refined grain size. The smaller the size of the microscopic spherical precipitates, the smaller the grain size after refinement. The spherical precipitated phase less than or equal to 2 µm has a significant effect on grain refinement. Acknowledgements The authors are pleased to acknowledge financial support received from the following projects (in no particular order). The National Key Research and Development Program of China (2016YFB0300804 and 2016YFB0300801), and the National Natural Science Fund for Innovative Research Groups (Grant No. 51621003). The Construction Project for National Engineering Laboratory for Industrial Big-data Application Technology (312000522303). National Natural Science Foundation of China (No. 51671005 and 51701006), Beijing Natural Science Foundation (2162006) and Program on Jiangsu Key Laboratory for Clad Materials (BM2014006).

References 1. T.E. Quested, Understanding mechanisms of grain refinement of aluminium alloys by inoculation. Mater. Sci. Technol. 20, 1357–1369 (2004) 2. L. Sturz, A. Drevermann, C. Pickmann, G. Zimmermann, Influence of grain refinement on the columnar-to-equiaxed transition in binary Al alloys. Mater. Sci. Eng. A, 413, 379 (2005) 3. Z. Gao, The world aluminum grain refiner supply industry (1). Light Met. (9), 50–53 (1999) 4. D.H. Stjohn, M. Qian, M.A. Easton et al., The interdependence theory: the relationship between grain formation and nucleant selection. Acta Mater. 59(12), 4907–4921 (2011) 5. B.S. Murty, S.A. Kori, M. Chakraborty, Grain refinement of aluminium and its alloys by heterogeneous nucleation and alloying. Int. Mater. Rev. 47, 3 (2002)

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6. Y.D. He, X.M. Zhang, Z.Q. Cao, Effect of minor Cr, Mn, Zr, Ti and B on grain refinement of as-cast Al-Zn-Mg-Cu alloys. Rare Met. Mater. Eng. 39, 1135 (2010) 7. Z. Fan, Y. Wang, Y. Zhang et al., Grain refining mechanism in the Al/Al–Ti–B system. Acta Mater. 84, 292–304 (2015) 8. Z.R. Nie, S.P. Wen, H. Hui et al., Research progress of Er-containing aluminum alloy. Chin. J. Nonferrous Met. 21(10), 2361–2370 (2011) 9. Z.-r. Nie, The effect and progress of alloying elements in aluminium. China Nonferrous Met. 22, 56–57 (2009) 10. Z.R. Nie, T.N. Jin, J.B. Fu, G.F. Xu, J.J. Yang, J.X. Zhou, T.Y. Zuo, Research on rare earth in aluminum. Mater. Sci. Forum 396(402), 1731–1736 (2002) 11. Z.R. Nie, B.L. Li, W. Wang, T.N. Jin, H. Huang, H.M. Li, J.X. Zou, T.Y. Zuo, Study on the erbium strengthened aluminum alloy. Mater. Sci. Forum 546(549), 623–628 (2007) 12. F. Wang, D. Qiu, Z.L. Liu et al., The grain refinement mechanism of cast aluminium by zirconium. Acta Mater. 61(15), 5636–5645 (2013) 13. S.P. Wen, K.Y. Gao, H. Huang, Z.R. Nie, Synergetic effect of Er and Zr on the precipitation hardening of Al-Er-Zr alloy. Scr. Mater. 65, 592–595 (2011) 14. S.P. Wen, Z.B. Xing, H. Huang, B.L. Li, W. Wang, Z.R. Nie, The effect of erbium on the microstructure and mechanical properties of Al-Mg-Mn-Zr alloy. Mater. Sci. Eng. A 516, 42–49 (2009) 15. H. Li, B. Jie, J. Liu et al., Precipitation evolution and coarsening resistance at 400 °C of Al microalloyed with Zr and Er. Scr. Mater. 67(1), 73–76 (2012) 16. Z.-b. Xin, Z.-r. Nie, J.-x. Zou, X.-d. Gao, Form and effect of element Er in Al-Er alloy cast ingot. J. Chin. Rare Earths Soc. 25(2), 234–238 (2007) 17. J.-j. Yang, Z.-r. Nie, T.-n. Jin, G.-f. Xu, J.-b. Fu, H.-q. Ruan, T.-y. Zuo, Effect of trace rare earth element Er on high pure Al. Trans. Nonferrous Met. Soc. China 13(5), 1035–1039 (2003) 18. G.-f. Xu, Z.-r. Nie, T.-n. Jin, J.-j. Yang, J.-b. Fu, Z.-m. Yin, Effects of trace erbium on casting microstructure of LF3 Al alloy. J. Chin. Rare Earth Soc. 20(2), 143–145 (2002) 19. B. Gong, S.-p. Wen, H. Huang, Z.-r. Nie, Evolution of nanoscale Al3(Er1−x Zrx ) precipitates in Al-6Mg-0.7Mn-0.1Zr-0.3Er alloy during annealing. Acta Metall. Sin. 46(7), 850–856 (2010)

Effect of In Situ TiB2 Particle Content on Microstructure and Properties of Cast Al–Si Alloy Hongda Wang, Lihua Chai and He Li

Abstract Cast Al–Si alloys are widely used for their excellent casting properties, but their applications are restricted in some special fields because of their low mechanical properties. In this paper, H3 BO3 and TiO2 were used as raw materials to prepare TiB2 /Al–Si composites with different mass fractions by in situ generation. The effects of TiB2 particle content on microstructure and properties of the alloy were studied. The results showed that the microstructure of the alloy was mainly composed of coarse primary α-Al dendrites and long-striped eutectic Si particles. The coarse primary α-Al particles were obviously refined and the edge of acicular eutectic Si was rounded with the addition of TiB2 particles. When the particle content was low, the particle distribution was dispersed, and a large number of particles were located in the grain boundary. When the particle content was high, the particles were agglomerated and gathered at the grain boundary. With the increase of particle content, the mechanical properties of the material were obviously enhanced. When the content of TiB2 particles was 5%, the tensile strength and yield strength reached 200 and 160 MPa, which were 22.8 and 34% higher than those of matrix, respectively.

1 Introduction Aluminum and its alloys have the characteristics of low density, high specific strength, good electrical conductivity, and good thermal conductivity. They are widely used in automotive, aerospace, weapon, and equipment manufacturing [1, 2]. Among them, the casting aluminum-silicon alloy has good fluidity and filling ability. It is a common material for manufacturing automobile wheel hub and cylinder. However, the mechanical properties of Al–Si alloys are poor, which limits their further application [3]. The mechanical properties of aluminum and aluminum alloys can be improved effectively by adding smaller reinforcement particles to aluminum and aluminum alloys under the condition of ensuring elongation. However, the poor wettability of H. Wang · L. Chai (B) · H. Li College of Materials Science and Engineering, Beijing University of Technology, Beijing 100124, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_24

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the interface between the reinforced particles and the matrix results in the limitation of application and development of aluminum matrix composites prepared by the traditional addition method. Particle-reinforced aluminum matrix composites prepared by in situ reaction method can effectively overcome as mentioned above [4–6]. At present, the research of particle-reinforced aluminum matrix composites is mainly focused on the preparation process. The addition of TiB2 particles to the cast Al–Si alloy can effectively reduce the grain size and improve the strength of the material. The average size of TiB2 particles prepared by in situ reaction method is 1 μm [7]. However, the influence of particle state, especially the content and morphology of particles on the material is still lack of systematic study. In this paper, TiB2 /Al–Si composites were prepared by self-propagating high-temperature synthesis and melt reaction. The effects of particle content on microstructure and properties of the composites were studied.

2 Materials and Methods A Al–10 wt% Si alloy was used as the base material of the experiment, and the raw material was pure aluminum (99.9%, mass fraction) Al–20 wt% Si master alloy. Al/TiB2 master alloy was prepared by using boric acid (H3 BO3 ), titanium oxide (TiO2 ), titanium powder and aluminum powder as raw materials. The raw material is mixed and dried in the mole ratio and pressed into a blank using a crucible furnace to heat aluminum ingots to 950 °C. The preform was put into molten aluminum to ignite, and the reaction was maintained by the reaction exothermic. After the reaction was finished, the pure phase Al/TiB2 master alloy was prepared by slagging, refining and casting. Al/TiB2 master alloy was remelted at 750 °C and TiB2 /Al–Si composites with different particle content were prepared according to the content of Al–20 Si alloy. A ϕ100 mm × 80 mm ingot was obtained by pouring the melt into a metal mold with preheating temperature of 200 °C. The microstructure was observed by sampling in the middle of the ingot, and the tensile specimen of ϕ5 mm × 25 mm was processed. The optical microstructure of the composite was observed by OLYMPUS PMG3 optical microscope after grinding and polishing, the phase constitution was analyzed by D/MAX-3C rotating anode X-ray diffractometer, and the microstructure was analyzed by QUANTA200 scanning electron microscope. A series of tensile tests were carried out on Instron5569 electronic universal material testing machine. The tensile rate was 1 mm/min and the average value of mechanical properties was obtained.

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3 Experimental Results and Discussion 3.1 Analysis of Master Alloy Al/TiB2 master alloys were analyzed by SEM, XRD, EDS method. Figure 1a and b show the microstructure of Al/TiB2 master alloy. It can be seen that TiB2 particle distribution is relatively uniform on the matrix with hexagonal and square shape, and the average size of TiB2 particles is 1.1 μm. Due to the large surface tension of the particles, particle aggregation is inevitable in a small area. Figure 1c and d is the result of EDS and XRD analysis of Al/TiB2 master alloy. It can be seen from the spectrum that only Al and TiB2 phases exist in the master alloy, and no brittle hard TiAl3 phase Presents which shows that the pure Al/TiB2 master alloy can be prepared by this method.

(a)

(b)

(c)

(d)

Fig. 1 Microstructure of Al/TiB2 master alloy. a Microstructure (low magnification), b microstructure (high magnification), c EDS analysis and d XRD analysis

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3.2 Microstructure of Composite Materials Figure 2 shows the composite microstructure with different particle content (0, 1, 3 and 5 wt%). As shown in Fig. 2a, the matrix alloy mainly consists of dendritic primary α-Al phase and acicular eutectic Si. The white region is the primary α-Al phase and the grey precipitates in coarse dendritic formed during solidification. Some black needle-like structure distributes at the adjacent parts of primary α-Al, which is Sienriched phase in eutectic structure. With the addition of TiB2 particles, the primary α-Al changes from coarse dendritic to rose-shaped, and the existence of primary α-Al particles increased the growth resistance of α-Al, hindered the grain growth and dendritic formation, and made the needle-shaped Si edges rounded. When the particle content is low, the particles distribute dispersedly in the matrix, and with the increase of the particle content, a small range of segregation occurs at the grain boundary, as shown in Fig. 3. The reason is that during solidification process, it is difficult for the particles to be swallowed and moved to the grain boundary with the

Fig. 2 Composite microstructure under different particle content a 0 wt%, b 1 wt%, c 3 wt% and d 5 wt%

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Fig. 3 TiB2 particles distribution under different particle content a 0 wt%, b 1 wt%, c 3 wt% and d 5 wt%

migration of the solid–liquid interface, and it is easy to reduce the surface energy by agglomeration under casting conditions, so the phenomenon of local agglomeration occurs.

3.3 Mechanical Properties of Composites Figure 4 shows the Brinell hardness. It can be seen from the chart that the hardness of the composites increases with the increase of particle content. When the particle content is 5%, the hardness of the composites is 41.7 HB, which is 8.3% higher than that of the matrix. Table 1 shows the tensile properties of composites with different particle contents (0, 1, 3, and 5 wt%). It can be seen from the table that with the increase of TiB2 particle content, the mechanical properties of the composites are improved gradually. When the particle content is 5%, the tensile strength and yield

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Fig. 4 Brinell hardness of the composites with different particle content (0, 1, 3 and 5 wt%)

Table 1 Tensile properties of the composites with different particle content

TiB2 particle content (wt%)

0

1

3

5

Tensile strength (MPa)

166.82

169.79

188.48

205.78

Yield strength (MPa)

110.51

125.97

139.42

154.77

4.94

2.83

3.07

2.55

Elongation (%)

strength of the composites reach 205.78 and 154.77 MPa respectively, compared with the matrix, the tensile strength and yield strength of the composites are increased by 23.4 and 40.1%, respectively. The addition of TiB2 particles leads to a decrease in the elongation of the composites. The reason is that stress concentration occurs near some large particle and leads to cracking under stress state. As the particles increased, the elongation rate deceased. For composites, the strengthening effect mainly refers to the improvement of tensile strength and yield strength. The main causes of composite strengthening are the synergistic effect of load transfer mechanism, Orowan reinforcement mechanism and grain refinement mechanism [8]. When the composite material receives the external force, the matrix and the reinforced particle bear the load together. Because of the coordinated deformation, the stress of the matrix is small, and the strength and the bearing capacity of composite are improved. When the TiB2 particles are small dispersive particles, the reinforcement of the material will also be enhanced by the dislocation movement caused by the addition of reinforcing phase, thus forming the Orowan strengthening [9, 10]. For the composite materials, the smaller TiB2 particles are difficult to be shearing because of their high strength, and the dislocation lines bypass the particles by the Orowan mechanism. The combination of matrix strength and particle will further enhance the strength of composites. The small-sized TiB2 particles prevent grain growth, produce grain boundary pinning, and greatly reduce

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Fig. 5 Fractographies of TiB2 /Al–Si composites

the stress concentration inside the composite material. The mechanical properties of the composites are greatly improved by the combined action of many factors. Figure 5 shows the tensile fracture morphologies of the composite, the fracture surface is flat, perpendicular to the tensile axis, indicating that the material is brittle fracture. The EDS analysis of the yellow line area for casting defects shows that the inclusions are Al2 O3 . It can be seen from the morphology of high magnification fracture that there are many silicon fracture surfaces and tearing ridges, showing the characteristics of quasi-cleavage fracture. The cracks mainly originate from the fracture of Si phase and the bond between Si phase and matrix. Due to the large stress concentration near Si phase, the brittle cracking of Si phase occurs. There are a large number of submicron TiB2 particles at the bottom of the dimple, which indicates that the TiB2 particles are mainly distributed at the interface between the particles and the matrix. This is due to the smaller size and higher strength of the TiB2 particles with a regular shape, which is not easy to crack due to the stress concentration. However, due to the difference of thermal expansion coefficient between the matrix and the particles, the interface between the reinforced particles and the matrix is in a loaded state. Finally, under the action of the external force, the particles appeared the phenomenon of destaining.

4 Conclusions The results are listed as follows: (1) TiB2 /Al–Si composites without brittle and hard TiAl3 phase can be prepared by a new in situ reaction method, and the TiB2 particle content is high. (2) The grain size of TiB2 particles is refined, and the primary Al changes from coarse dendrite to rose shape, and the edge of the eutectic Si is rounded.

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(3) When the particle content is low, the distribution in the matrix is more uniform. When the particle content is high, a local area of agglomeration occurs at the grain boundaries. With the increase of TiB2 particle content, the mechanical properties of the composite are improved. When the content of TiB2 particles is 5%, the tensile strength and yield strength reached 200 and 160 MPa, which are 22.8 and 34.8% higher than those of matrix, respectively.

References 1. J. Sun, X. Zhang, Y. Zhang, Effect of alloy elements on the morphology transformation of TiB2 particle in Al matrix. Micron 70, 21–25 (2015) 2. A. Rabiezadeh, A.M. Hadian, A. Ataie, Synthesis and sintering of TiB2 nanoparticles. Ceram. Int. 40(10), 15775–15782 (2014) 3. V.I. Nikitin, J.I.E. Wanqi, E.G. Kandalova et al., Preparation of Al–Ti–B grain refiner by SHS technology. Scripta Mater. 42(6), 561–566 (2000) 4. E.M. Sharifi, F. Karimzadeh, M.H. Enayati, Synthesis of titanium diboride reinforced alumina matrix nanocomposite by mechanochemical reaction of Al–TiO2 –B2 O3 . J. Alloy. Compd. 502(2), 508–512 (2010) 5. S. Lakshmi, L. Lu, M. Gupta, In situ preparation of TiB2 reinforced Al based composites. J. Mater. Process. Technol. 73(1), 160–166 (1998) 6. H. Sun, X. Li, D. Chen, H. Wang, Enhanced corrosion resistance of discontinuous anodic film on in situ TiB2 /A356 composite by cerium electrolysis treatment. J. Mater. Sci. 44(3), 786–793 (2008) 7. E. Wang, T. Gao, J. Nie, X. Liu, Grain refinement limit and mechanical properties of 6063 alloy inoculated by Al–Ti–C (B) master alloys. J. Alloy. Compd. 594, 7–11 (2014) 8. H. Yi, N. Ma, Y. Zhang, X. Li, H. Wang, Effective elastic moduli of Al–Si composites reinforced in situ with TiB2 particles. Scripta Mater. 54(6), 1093–1097 (2006) 9. X. Wang, Z. Liu, W. Dai, Q. Han, On the understanding of aluminum grain refinement by Al–Ti–B type master alloys. Metall. Mater. Trans. B 46(4), 1620–1625 (2014) 10. S. Amirkhanlou, M. Rahimian, M. Ketabchi, N. Parvin, P. Yaghinali, F. Carreño, Strengthening mechanisms in nanostructured Al/SiCp composite manufactured by accumulative press bonding. Metall. Mater. Trans. A 47(10), 5136–5145 (2016)

The Influence of Different Casting Methods on the Corrosion Resistance of Mg–Gd–Y Alloys Renju Cheng, Hong Tang, Wenjun Liu, Na Zhang, Hanwu Dong, Bin Jiang and Fusheng Pan

Abstract The present work mainly studied the corrosion resistance and corrosion process of Mg–Gd–Y alloys casted by low-pressure sand casting and gravity diecasting methods. The results showed that the main corrosion type of the Mg–Gd–Y alloys was galvanic corrosion, which were formed by α-Mg matrix with Mg3 (Gd, Y) or, Mg5 (Gd, Y) particle phases. The particle phases in the low pressure sand casting alloys were mainly Mg5 (Gd, Y) phases, but Mg3 (Gd, Y) in the gravity die-casting alloys. The Mg5 (Gd, Y) phases dissolved into α-Mg matrix by solid solution treatment (T4), but due to their high melting points the Mg3 (Gd, Y) phases could not dissolve efficiently. So after T4 treatment, the particles in the low-pressure sand casting alloy decreased and then its corrosion resistance was greatly improved. However, there were still lots of residual particles in the gravity die-casting alloys after T4 treatment, so their corrosion resistance did not change obviously.

1 Introduction As one of the lightest structural materials, magnesium alloys are widely used in the automotive and aerospace industries for their low density, high specific strength, superior damping capacity and good cast-ability [1–3]. In addition, magnesium alloy has good recyclability and can be reused for nearly 100%, so as known as “the green engineering material in twenty-first Century” [4, 5]. Unfortunately, magnesium has a high chemical reactivity that would easily induce the occurrence of corrosion in environment, its standard electrode potential is only −2.36 V [6]. The formed oxide R. Cheng (B) · W. Liu · N. Zhang · H. Dong · F. Pan Chongqing Academy of Science and Technology, Chongqing 401123, China e-mail: [email protected] H. Tang · B. Jiang · F. Pan College of Materials Science and Engineering, Chongqing University, Chongqing 400044, China B. Jiang · F. Pan National Engineering Research Center for Magnesium Alloy, Chongqing University, Chongqing 400044, China © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_25

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film is porous with density coefficient of 0.79, which cannot effectively protect the matrix [7] or impede the corrosion. At present, the corrosion resistance of magnesium alloys has become one of the main factors which hindering the application of magnesium alloys. Rare earth elements have relatively large solid solubility in magnesium alloys, and the solid solubility drops sharply with the reducing of temperature, so rare earth has obvious solution and aging strengthening abilities in magnesium alloys. The addition of rare earth elements can effectively improve the organization, increase the strength and corrosion resistance of magnesium alloys [8]. Mg–Nd, Mg–Y, and Mg–Gd alloys are the most typical rare earth magnesium alloys. Among them, Mg–Gd–Y alloys have attracted wide attention because of their good room temperature strength and high-temperature creep resistance [9]. But the current researches are mainly based on mechanical properties, and there are few studies on the corrosion characteristics, especially the influence of casting methods on the corrosion resistance of the Mg–Gd–Y alloys.

2 Experimental Procedures The alloys were prepared from commercial pure Mg, pure Ag, Mg–30 wt% Gd, Mg–30 wt% Y, and Mg–30 wt% Zr master alloys. Melting was carried out in a stainless steel crucible placed in an electric resistance furnace under a mixed atmosphere of CO2 and SF6 with the ratio of 6:1. The ingots of gravity die-casting were cast in a steel mold pre-heated up to about 350 °C. The casting process of low pressure sand casting and gravity die-casting are shown in Fig. 1, the low-pressure sand casting melting metal was reversed into sand mold pre-heated to about 350 °C by controlling gas and holding pressure for 10 min and for gravity die-casting the melting metal was poured into the mold pre-heated to about 350 °C. Specimens were solution treated at 500 °C for 10 h then quenched into water to room temperature (T4 condition), and then isothermal aged at 225 °C for 36 h (T6 condition).

Fig. 1 Schematic diagram of the casting methods: a low pressure sand casting; b gravity die-casting

The Influence of Different Casting Methods on the Corrosion … Table 1 Chemical composition of the alloys (wt%)

Casting method

Mg

Low pressure sand casting Gravity die-casting

251

Gd

Y

Ag

Zr

89.06 7.83

1.46

1.13

0.43

89.12 7.76

1.47

1.20

0.45

The SEM and OM samples of the alloys were prepared by electron polish with the solution of 5 vol% perchloric acid in ethanol. The hydrogen evolution experiment was carried out at room temperature with PH value of 6.6 by using 3.5 wt% NaCl solution as corrosion liquid. The SEM observations were performed on a TESCAN VEGA 2 LMU. The compositions of the castings were analyzed by an inductively coupled plasma atomic emission spectrometry (ICP-AES), and the exact compositions were listed in Table 1.

3 Results and Discussion Figure 2 shows the hydrogen evolution results of the alloys casted by both low pressure sand casting (L) and gravity die-casting (G) methods with the states of as-casted (F), solution treated (T4) and aging treatment (T6). The corresponding hydrogen evolution rates are listed in Table 2. It can be seen that the hydrogen evolution rates of the low pressure sand casting alloys are higher than that of the gravity die-casting alloys both at the state of as-cast and T6, however after T4 treatment, the hydrogen

Fig. 2 Corrosion experiment results of the alloys

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Table 2 Hydrogen evolution rates for the alloys Casting method

The rate of hydrogen evolution (ml/(cm2 h)) As-cast

T4

T6

Low pressure sand casting

0.42

0.09

0.84

Gravity die-casting

0.11

0.14

0.40

Fig. 3 SEM microstructures of the as-casted alloys: a low pressure casting; b gravity die-casting Table 3 The EDS results of the as-casted alloys (at%)

Element

Mg

Gd

Y

Ag

Zr

A

86.11

9.66

2.28

B

58.17

23.23

18.60

C

87.92

8.24

2.04

D

85.45

7.57

3.89

E

70.07

10.85

19.08

F

63.11

15.36

21.52

–

–

1.96

–

–

–

1.79

–

3.09

–

–

–

evolution rate of the low pressure sand casting alloys greatly reduced, but increased for the gravity die-casting alloys. Figure 3 and Table 3 show the typical microstructures and its corresponding EDS analysis results of the as-casted alloys. The XRD results (Fig. 4) show that the second phases are constituted by island phases, particle phases and square phases. Combine with the previous studies, we can infer that the island phases are Mg24 (Gd, Y)5 phases, the particle phases are Mg5 (Gd, Y) and Mg3 (Gd, Y) phases, the square phases are the Mg–Gd–Y compound [10]. Figure 5 and Table 4 show the typical microstructures and its corresponding EDS analysis results of the alloys at T4 state, combine with the XRD results shown in Fig. 6, we can see that the island Mg24 (Gd, Y)5 phases have been mostly dissolved

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Fig. 4 The XRD results of the as-casted alloys: a low pressure casting; b gravity die-casting

Fig. 5 SEM microstructures of the T4 treated alloys: a low pressure casting; b gravity die-casting

into the matrix after T4 treated, only particle and square phases remain in the matrix. Meanwhile, compared with Figs. 4 and 6, we can see that the diffraction peaks of Mg5 (Gd, Y) phases disappeared after T4 treated, and the peak intensity of Mg24 (Gd, Y)5 phases decreases. According to the results of EDS spectrum analysis (in Table 4), the particle phases in the alloy after T4 treated are mainly composed of Mg3 (Gd, Y) phases, and the number of residual particle phases in the gravity die-casting alloy is much more than that of it in the low pressure sand casting alloys. Figure 7 shows the morphology of the as-casted alloys after immersion in the 5 wt% NaCl solution for 60 min, from which we can see that the corrosion areas are

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Table 4 The EDS results of the T4 treated alloys (at%) Element

Mg

Gd

Y

Ag

Zr

A

40.24

31.24

28.52

–

–

B

69.70

16.59

13.71

–

–

C

30.68

39.49

29.83

–

–

D

63.11

15.36

21.52

–

–

Fig. 6 The XRD results of the T4 treated alloys: a low pressure casting; b gravity die-casting

mainly concentrated around the particle phases. It is deduced that the corrosion is mainly due to the galvanic corrosion formed by the α-Mg matrix with the particle Mg3 (Gd, Y) or Mg5 (Gd, Y) phases [11, 12]. The matrix around the particle phase is then continuously dissolved, resulting in the weakening of the binding force between the second phases and the matrix, and then caused the fall off of the second phase particles. The solidification rate of the alloys are pretty slowly during the solidification process of the low pressure sand casting alloys, so the solidification process is very close to the equilibrium solidification process, make the solidified structure mainly composed of α-Mg matrix and eutectic structures. The Gd and Y elements in the alloy tend to form Mg5 (Gd, Y) and Mg24 (Gd, Y)5 phases. But the solidification rate of gravity die-casting alloys are much faster, during the solidification the composition deviates from the equilibrium solidification process. As the composition fluctuates, lots of Mg3 (Gd, Y) phases are formed by composition fluctuation and cannot be fully diffusing. As large number of Gd elements were consumed in advance, the amount of Mg3 (Gd, Y) and Mg24 (Gd, Y)5 phases formed in the gravity die-casting method is relatively small. Therefore, the particle phases in the low-pressure sand casting

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Fig. 7 SEM image of the as-cast alloy after immersion in 5 wt% NaCl aqueous for 60 min

alloys are mainly composed of Mg5 (Gd, Y) phases, while in the gravity die-casting alloy are mainly composed of Mg3 (Gd, Y) phases. The same content of Gd and Y elements can form more Mg5 (Gd, Y) phases than Mg3 (Gd, Y) phases, so the amount of particle phases in the low-pressure sand casting alloys is more than in the gravity die-casting alloys, so the corrosion area in the low-pressure sand casting alloy is larger than that of the gravity die-casting alloys, which leads to the higher hydrogen evolution rate of the low-pressure sand casting alloys than the gravity die-casting alloys. The Mg5 (Gd, Y) phases can dissolve into the α-Mg matrix after T4 treated, but due to its high melting point the Mg3 (Gd, Y) phases are hard to be dissolved into the α-Mg matrix. So after T4 is treated, the particles in the low-pressure sand casting alloy decreases and then its corrosion resistance is greatly improved. But there are still lots of residual particles in the gravity die-cast alloys after T4 is treated, so its corrosion resistance does not change a lot. However, the corrosion resistance of the alloys decreases again due to the precipitation of the second phase after the T6 is treated.

4 Conclusions 1. The hydrogen evolution rates of the low-pressure sand casting alloys are higher than those of the gravity die-casting alloys both at the state of as-cast and T6. However, after T4 treatment, the hydrogen evolution rate of the low-pressure sand casting alloys greatly reduces, but increases for the gravity die-casting alloys. 2. The corrosion of the Mg–Gd–Y alloys is mainly due to the galvanic corrosion formed by the α-Mg matrix with the particle Mg3 (Gd, Y) or Mg5 (Gd, Y) phases.

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3. The particle phases in the low-pressure sand casting alloys are mainly composed of Mg5 (Gd, Y), while in the gravity die-casting alloys are mainly composed of Mg3 (Gd, Y). The Mg5 (Gd, Y) phases can dissolve into the α-Mg matrix by solid solution treatment (T4), so the corrosion resistance is greatly improved after T4 treatment. But due to their high melting points, the Mg3 (Gd, Y) phases are very hard to dissolve into the α-Mg matrix, so after T4 treatment the corrosion resistance does not change obviously. Acknowledgements This work is supported by the National Key R&D Program of China (2016YFB0301100), National Natural Science Foundation of China (NSFC) (51504052), Chongqing Special key technology innovation of key industries (cstc2017zdcy-zdzxX0006) and Chongqing technological innovation and application demonstration (cstc2018jszx-cyzdx0082).

References 1. Z. Yang, J.P. Li, J.X. Zhang et al., Review on research and development of magnesium alloys. Acta Metall. Sin. 26(3), 217 (2013) 2. X. Liu, D. Shan, Y. Song, E-h Han, Influence of yttrium element on the corrosion behaviors of Mg–Y binary magnesium alloy. J. Magnes. Alloy. 5(1), 26–34 (2017) 3. A.A. Luo, Magnesium casting technology for structural applications. Magnes. Alloy. 1(2), 115 (2013) 4. Y. Wang, Z.D. Liao, H.L. Zhang et al., Development on effects of rare earth element in magnesium alloy. Mater. Rev. S1(25), 487–491 (2011) 5. W.B. Du, K. Liu, K. Ma, Z.H. Wang, S.B. Li, Effects of trace Ca/Sn addition on corrosion behaviors of biodegradable Mg–4Zn–0.2Mn alloy. J. Magnes. Alloy. 6(1), 1–14 (2018) 6. Z.H. Chen, H.G. Yan, J.H. Chen, Magnesium Alloy (Chemical Industry Press, Beijing, 2004) 7. I. Polmear, Magnesium alloys and application. Mater. Sci. Technol. 10, 1 (1994) 8. Q.W. Chen, A.T. Tang, T.X. Xu et al., High performance cast magnesium rare-earth alloys: retrospect and prospect. Mater. Rev. 30(9), 1–9 (2016) 9. H.R. Nodooshan, W.C. Liu, G.H. Wu et al., Effect of Gd content on microstructure and mechanical properties of Mg–Gd–Y–Zr alloys under peak-aged condition. Mater. Sci. Eng., A 615, 79–86 (2014) 10. Y. Gao, Q.D. Wang, J.H. Gu et al., Mechanical properties and creep behavior of Mg–Gd–Y alloys. Mater. Sci. Forum 546, 163–166 (2007) 11. G.Y. Zhang, H. Zhang, Z.F. Zhao, Electronic theoretical study of the influence of impurities on corrosion resistance of magnesium alloy. Acta Phys. Sin. 55(5), 2439–2443 (2006) 12. M.Q. Zhao, A.L. Lei, Corrosion and Protection of Metal (National Defense Industry Press, Beijing, 2014)

Effect of Cu Additions and Extrusion Treatment on the Microstructure and Mechanical Properties of Mg–6Sn–1Al Alloy Zhijian Ye, Tong Li, Gui Lou, Jianhang Yue and Xinying Teng

Abstract To further promote the industrial production and expand the applications of magnesium alloys in the field of lightweight structural materials, the researches of low-cost, rare-earth-free magnesium alloys have received extensive attention. Mg–Sn–Al alloys possess great potentials in the application. However, their mechanical properties are relatively poor, which is incapable of meeting the requirements of structural materials for replacing rare-earth magnesium alloy. In this paper, the microstructure of Mg–6Sn–1Al–xCu (x  0.5, 1, 2) (wt%) alloys were characterized by XRD and SEM. The mechanical properties of the alloys were investigated by combined tensile and compression tests. The results showed that the addition of Cu in Mg–6Sn–1Al alloys had a positive effect on mechanical properties, which was mainly attributed to the fine-grained strengthening and second-phase strengthening. In addition, comparing with as-cast alloys, extrusion processing had led to significant changes in the microstructure of the alloy. Due to the recrystallization behaviors during the hot extrusion process, the eutectic structure of Mg2 Sn (Al, Cu) disappeared and transformed into granular precipitates dispersing on the grain boundary. From the tensile and compressive test, it was found that the extruded Mg–6Sn–1Al–0.5Cu alloy exhibited the highest ultimate tensile strength (UTS) of 300.1 MPa, which was 131% higher than that of the cast alloy. With the increase of Cu addition, the 0.2% compressive yield strength increased from 72 MPa of Mg–6Sn–1Al–0.5Cu alloy to 110 MPa of Mg–6Sn–1Al–2Cu alloy, but their ultimate compressive strength (UCS) remained stable.

1 Introduction In recent years, magnesium alloys with high specific strength have been used as structural components in automobile, aerospace and portable consumer electronic device applications [1, 2]. However, restricted strength, ductility, and processing formabilZ. Ye · T. Li · G. Lou · J. Yue · X. Teng (B) School of Materials Science and Engineering, University of Jinan, No. 336, West Road of Nan Xinzhuang, Jinan 250022, People’s Republic of China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_26

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ity have limited their wide applications. Therefore, many researchers focused on the strengthening design of Mg alloys and have achieved the desired results. Alloying elements addition and deformation process are effective ways to improve the mechanical properties of Mg alloys. Teng et al. [3] reported Mg–Zn–Y quasicrystalreinforced AZ91 alloys, which indicated that after adding 6 wt% Mg30 Zn60 Y10 quasicrystal master alloy, the ultimate tensile strength (UTS) of the Mg-based quasicrystal composites could reach the peak, 214.85 MPa, which increased about 33% than that of AZ91 matrix alloy. Yu et al. [4] developed a high strength Mg–11Gd–4.5Y–1Nd–1.5Zn–0.5Zr (wt%) alloy with a tensile yield strength (TYS) of 482 MPa, an UTS of 517 MPa and an elongation to failure of 2% at room temperature through the processing methods of hot extrusion, cold rolling and aging treatment. Though Mg-RE alloys exhibit high strength and toughness as well as favorable processing formability, the high cost and resource scarcity of rare-earth elements restrict their varied applications to a certain extent. In this content, importance has been increasingly attached to researches into rareearth-free Mg alloys by a large majority of investigators. Among them, Mg–Sn-based alloy is one of the most potential alloy systems for developing both casting and wrought alloy products [5]. In Mg–Sn binary system, Mg2 Sn has been proved to be a thermally stable phase with the melting point of 770 °C [6] higher than that of Mg17 Al12 phase (420 °C) in Mg–Al-based alloys. According to the Mg–Sn binary phase diagram, the solid solubility of Sn in α-Mg is 14.5 wt% at the eutectic temperature (561 °C), and it decreases to 0.45 wt% at 200 °C, which provides a large variation for precipitation strengthening [7]. However, in Mg–Sn alloys, Mg2 Sn phase tends to appear as slender rods or coarse particles, which produce an undesirable influence on the optimization of properties. Consequently, various researches have been carried out in Mg–Sn alloys via alloying elements additions such as Zn, Al, Ca, Zr, Si, etc. [8–10]. Among these alloying elements, Al and Cu are promising. Al is one of the most popular and widely commercialized alloying elements. A large number of β-Mg17 Al12 phases are uniformly dispersed in the matrix and grain boundaries, which enhance the resistance to deformation and destruction [11]. The addition of Cu into Mg results in Mg2 Cu intermetallic phase. Due to the limited solubility of Cu in Mg, Mg2 Cu phases are generally present along the grain boundaries [12]. Although the addition of Cu could improve the strength and ductility of alloys to some extent [13], it reduces the corrosion resistance [14]. On the basis of previous investigations, we are committed to studying the effects of Cu alloying on microstructure and mechanical properties of as-cast and extruded Mg–Sn–Al alloys. Therefore, new alloys with the detailed composition of Mg–6Sn–1Al–xCu (x  0.5, 1, 2) (wt%) are designed and prepared. The aim of our research is to develop novel rare-earth-free alloys that have similar performance to Mg-RE alloy.

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2 Experimental Procedure The experimental alloys were prepared by melting at 720 °C in high-temperature resistance furnace using graphite clay crucibles under CO2 + SF6 mixture atmosphere, and then casting into iron molds preheated to 200 °C. Raw materials used for melting were high-pure Mg (99.99%), Sn (99.99%), Al (99.99%), and Mg–30Cu (wt%) master alloy. The actual compositions of the alloys were determined by area scan in SEM and EDS analyses and presented in Table 1. Before casting, the melt was preserved at 720 °C for 30 min to ensure composition uniformity. Alloy ingots were homogenized at 360 °C for 20 h before extrusion processing. After homogenization, the ingots were hot extruded into bars of 21 mm in diameter at 400 °C using an XJ630 Horizontal Extrusion Machine. The extrusion ratio was 23:1 and extrusion rate was set at 2 m/min. Extruded bars were processed into tensile testing rods and compressive cylinders. Tensile tests and compressive tests of as-cast and extruded alloys were carried out on WDW-100A microcomputer control electronic universal testing machines at a loading rate of 1 mm·min−1 . Repeated tests were performed at least three times in order to ensure the independence and accuracy of experimental results. The microstructure of Mg–6Sn–1Al–xCu (x  0.5, 1, 2) (wt%) were characterized by optical microscope (OM) and scanning electron microscopy (SEM) equipped with an energy dispersive spectroscopy (EDS). OM and SEM specimens were polished by 1 μm magnesium oxide powder suspension. Phase analysis was conducted on X-ray diffraction (XRD) (Cu Kα) with a scanning rate of ~5° min−1 . The average grain size was determined by using the Image J analysis software.

Table 1 The chemical compositions of experimental alloys Alloy Design composition (wt%)

Actual composition (wt%)

Actual composition (at%)

Sn

Al

Cu

Mg

Sn

Al

Cu

Mg

TAC- Mg–6Sn–1Al–0.5Cu 1

6.09

0.91

0.42

Bal.

1.32

0.86

0.17

Bal.

TAC- Mg–6Sn–1Al–1Cu 2

6.21

0.87

1.20

Bal.

1.35

0.83

0.49

Bal.

TAC- Mg–6Sn–1Al–2Cu 3

5.91

0.90

1.97

Bal.

1.29

0.86

0.80

Bal.

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3 Results and Discussion 3.1 Microstructure Figure 1 shows the XRD patterns of as-cast and as-extruded Mg–6Sn–1Al–xCu (x  0.5, 1, 2) alloys. It is found that the microstructure of the alloy mainly contains two structures: α-Mg and Mg2 Sn phases. As shown in Fig. 1a, when the amount of Cu element added in the as-cast alloy is 0.5 and 1 wt%, no diffraction peak of Mg2 Sn phases is detected in the sample due to the small volume fraction of Mg2 Sn phases. Further observation of the second-phase microstructure by OM and SEM shows that the volume fraction of the precipitated phases at the grain boundary is relatively low, and therefore the corresponding diffraction peaks are not shown in the XRD pattern

Fig. 1 XRD patterns of Mg–6Sn–1Al–xCu (x  0.5, 1, 2) (wt%) alloys: a as-cast; b extruded

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which is consistent with the results of the OM analysis. While adding 2 wt% Cu, the diffraction peaks of Mg2 Sn phases are detected in XRD, and Mg2 Sn phases are also observed in the microstructures of corresponding alloys. Thus, the addition of Cu promotes the precipitation of Mg2 Sn phases and the volume fraction of the second phase increases continuously. According to the phase diagram [6], it is concluded that the content point of Sn in the Mg–6Sn–1Al–xCu (x  0.5, 1, 2) alloys does not fall into the range of eutectic composition. Therefore, after adding Cu, the solid solubility of Sn in Mg is reduced. The left end of the eutectic composition section moves to the left. When the amount of Cu added reaches 2.0 wt%, Sn component corresponding to the left end of the actual eutectic section is less than 6 wt%. Therefore, pseudo-eutectic Mg2 Sn phases occur in the Mg–6Sn–1Al–2Cu alloy, corresponding to the presence of the Mg2 Sn diffraction peaks detected in the XRD patterns. Contrasting the as-cast alloy with the as-extruded alloy on XRD patterns, as shown in Fig. 1a and b, it is found that the relative intensity of the α-Mg diffraction peaks identified in the XRD is changed, which mainly due to the transformation of the grain orientation during hot extrusion process [15]. Since magnesium alloy is the hexagonal metal, basal slip is the most easily activated slip system during plastic deformation. During the hot extrusion process, the recrystallized grains rotate and are distributed along the basal slip direction. The orientation of α-Mg grains on the cross section of the extruded alloy bar (perpendicular to ED) is mostly parallel to the prismatic plane (10–10), while the orientation of α-Mg grains on the longitudinal section (parallel to ED) are mostly parallel to the basal plane (0002). (XRD samples of as-cast and extruded alloys in this experiment were both taken on the cross section of the extruded alloy bars.) Figure 2 shows OM images of the as-cast Mg–6Sn–1Al–xCu (x  0.5, 1, 2) alloys. It can be observed that with the increase of Cu addition, the microstructure morphology of the as-cast alloy changes significantly. As shown in Fig. 2a, the microstructure of as-cast alloy is mainly composed of α-Mg and α-Mg + Mg2 Sn eutectic structures. The α-Mg matrix consists of a majority of equiaxed grains and a few coarse columnar/dendrites grains at the same time, while the eutectic structure is distributed on the grain boundary of α-Mg. When Cu addition is increased to 0.5 wt%, Mg2 Sn phases possess a semicontinuous reticulate distribution on the grain boundaries with a low volume fraction; due to the low solid solubility of Cu in the Mg matrix as well as few addition of Cu, there are no independent Mg2 Cu phases formed during the solidification. Therefore, the Cu element tends to segregate on the grain boundary or in the eutectic structure. The solubility of Al in Mg is relatively high, so when Al is added with a small amount (1 wt%), Al atoms are almost dissolved into the Mg matrix, and only a small amount appears on the grain boundaries forming the Mg17 Al12 phase and coexisting with the Mg2 Sn. The study of Nan [16] showed that in the microstructure of as-cast Mg–3Al–3Sn alloy, a large number of continuous reticulate Mg17 Al12 and Mg2 Sn phases are distributed on the coarse grain boundary, while Mg17 Al12 phase is attached to Mg2 Sn phase. With the increase of Cu addition, the volume fraction of precipitate shows signs of steady rises. To compare the microstructures between alloys containing 0.5 and 1.0 wt% Cu, as

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(a)

(b)

(c)

(d)

(e)

(f)

Fig. 2 OM images and grain size distribution of as-cast Mg–6Sn–1Al–xCu (wt%) alloys: a, b x  0.5; c, d x  1; e, f x  2

Effect of Cu Additions and Extrusion Treatment … Table 2 EDS analysis results of positions in Fig. 4 (at%)

263

Alloy

Position

Mg

Extruded TAC-2

A

89.81

B

89.18

C

98.44

0.14

D

79.83

20.17

E

91.69

F

98.34

As-cast TAC-3

Sn

Al

Cu

0.11

5.44

4.64

10.12

0.70

0

1.01

0.41

0

0

6.45

1.55

0.32

0.54

1.12

0

shown in Fig. 2a and c, it can be found that the slight increase of the Cu addition makes the refinement of α-Mg grains, the total length of grain boundaries rises, and the volume fraction of the Mg2 Sn phase increases. When Cu addition reaches 2 wt%, as shown in Fig. 2e, a large area of the eutectic structure and pre-eutectic Mg2 Sn phases appeared. It can be clearly observed that the lamellar eutectic structure of Mg2 Sn (Al, Cu) (blue circle) and the coarse fish-bone or lath-shaped Mg2Sn phases (red circle) form at the grain boundary. Figure 4b is a backscatter SEM image of the as-cast Mg–6Sn–1Al–2Cu alloy with bright white fish-bone or lath-shaped second phases and lamellar eutectic structure distributed on the black α-Mg matrix. The EDS analysis (Table 2) reveals that the second phases are basically Mg2 Sn phases, and the eutectic structure mainly contains Mg2 Sn and Al–Cu intermetallic compounds. Figure 2b, d, and f reveal the grain size distribution of Mg–6Sn–1Al–xCu (x  0.5, 1, 2) alloys, respectively. It can be observed that the proper addition of Cu could refine α-Mg grains and optimize the microstructure of the alloy. When adding 0.5 wt% Cu, the average grain size is 35.4 μm; with Cu addition increased to 1.0 wt%, average grain size reaches to 24.2 μm, which is 30% reduction to the former. When the addition reaches to 2.0 wt%, a quantity of α-Mg + Mg2 Sn eutectic structure appears on the grain boundary. Due to constitutional supercooling caused by the segregation of alloying element Al, Cu in the eutectic structure, the solidification, and growth of α-Mg grains tend to be dendritic and the average grain size increases to 42.5 μm. There are two main reasons for grain refinement of as-cast alloys. First of all, the segregation of Cu elements on the grain boundaries causes the formation of concentration gradient at the solid–liquid interface during solidification and crystallization, which resulting in constitutional supercooling. In this condition, the crystalline cores before solid–liquid interface will be activated, thereby increasing the nucleation rate and leading to grain refinement [17, 18]. Second, the addition of Cu decreases the solid solubility of Sn in Mg, promoting the precipitation of Mg2 Sn phases and increasing the volume fraction of the second phase on the grain boundaries. The second phases on the grain boundaries suppress the growth of the α-Mg grains during solidification, thereby refining grains [19]. Figure 3 shows OM images of extruded Mg–6Sn–1Al–xCu (x  0.5, 1, 2) alloy. Due to the dynamic recovery during hot extrusion and the recrystallization during subsequent cooling process, the small amount of coarse α-Mg columnar grains in the as-cast alloy disappeared and were replaced by the regular equiaxed, recrystallized

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(a)

(b)

(c)

(d)

(e)

(f)

Fig. 3 OM images and grain size distribution of extruded Mg–6Sn–1Al–xCu (wt%) alloys: a, b x  0.5; c, d x  1; e, f x  2

grains; coarse precipitates, solid solution zone, and eutectic structure dissolved and second phases were precipitated in the grain boundaries and the inner grains. As shown in Fig. 3a and b, the microstructure of extruded Mg–6Sn–1Al–0.5Cu alloy is equiaxed α-Mg grains, with an average grain size of 39.47 ± 3.1 μm, and second phase particles had black spherical shape with an average size of 1.80 ± 0.22 μm, most of which are distributed on the grain boundaries, also a small amount existing inside the grains. When Cu addition reaches 1.0 wt%, as shown in Fig. 3c and d,

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the average grain size of the equiaxed α-Mg grains is 33.11 ± 2.7 μm. The number of second phase particles increases significantly with an average size of 3.39 ± 0.32 μm, which has more diffusion distribution on the matrix. With the Cu addition increased to 2.0 wt%, as shown in Fig. 3e and f, the average grain size of equiaxed grains is reduced to 29.59 ± 2.1 μm and the average grain size of the second phase particles is decreased to 2.57 ± 0.40 μm, but the amount is approximately the same as the former alloy, and a great deal of particles tend to be wire-like, leading to local agglomeration.

Fig. 4 SEM images of Mg–6Sn–1Al–xCu (wt%) alloys: a extruded, x  1; b as-cast, x  2

(a) B

C

A

(b)

F E

D

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Figure 4a shows a secondary-electron SEM image of extruded Mg–6Sn–1Al–1Cu alloy. Dispersed second phase particles are distributed on the equiaxed matrix. EDS analysis (as shown in Table 2) indicates that the second phases with larger grain size are AlCu phases, while the other smaller particles are Mg2 Sn phases. After hot extrusion, Sn has almost no solid solution in the matrix. Most of them form precipitates, and only a small amount exist on the grain boundary. It was reported that in Mg–Al–Sn alloys, the Mg17 Al12 phases formed by the accumulation of aluminum on the grain boundaries will adhere to the Mg2 Sn phases [16]. Therefore, during the hot extrusion process, Mg2 Sn phases precipitated in the first place and then coated with the precipitates formed by Al and Cu, so the amount of single Mg2 Sn phases observed and detected was small. Grain refinement of extruded alloys is mainly because of the recrystallization after hot extrusion. As the deformation process begins and proceeds, the density of dislocation continues increasing resulting dislocation tangle to form cellular substructures. At the same time, the increase of dislocation density also leads to the dynamic recovery process, and the disappearance rate of dislocation also rises up. Therefore, the proliferation rate and disappearance rate of dislocations are equal and achieve a stable state. At this time, dislocations are concentrated on the cell walls and form sub-grains, which keep an equiaxed state during deformation. By controlling the parameters of hot-working reasonably, a smaller size of the initial sub-grain structure can be obtained with lower deformation temperature and higher strain rate, providing a beneficial foundation for the recrystallized grains after cooling. In addition, since a large amount of second-phase particles precipitate during the hot extrusion process, grain boundaries are pinned and the movement of the grain boundary is hindered, thereby suppressing grain growth during the cooling process and obtaining finer grains.

3.2 Mechanical Properties The tensile properties of as-cast Mg–6Sn–1Al–xCu (x  0.5, 1, 2) alloys are shown in Fig. 5a and Table 3. With the increasing amount of Cu addition, tensile properties of as-cast alloys increase first and then decrease. The UTS of the as-cast Mg–6Sn–1Al–0.5Cu alloy is 130.0 ± 1.7 MPa and the elongation is 13.4%. When the Cu addition reaches 1 wt%, UTS is the highest, 166.7 ± 2.3 MPa, which increase by 28.3% over the former; its elongation increases to 14.3% as well. When the addition of Cu continues to rise up to 2.0 wt%, the UTS and elongation of the as-cast alloy decrease to 118.0 ± 2.7 MPa and 13.9%, respectively. By comparing the microstructure of the as-cast alloys, it is found that the grain refinement resulting from the addition of Cu is the main reason for the strengthening of the mechanical properties. During the process of plastic deformation, a sufficient number of dislocations must be accumulated inside the grains to provide the necessary stress, so that dislocation sources in adjacent grains start to move and produce macroscopic plastic deformation. Grain refinement can significantly increase the

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Fig. 5 Tensile engineering stress–strain curves of Mg–6Sn–1Al–xCu (x  0.5, 1, 2) (wt%) alloys: a as-cast; b extruded Table 3 Tensile mechanical properties of as-cast Mg–6Sn–1Al–xCu (x  0.5, 1, 2) (wt%) alloys Alloy

Ultimate tensile strength, UTS (MPa)

0.2% tensile yield strength, TYS (MPa)

Fracture strain (%)

As-cast TAC-1

130.0 ± 1.7

62.1 ± 2.1

13.4

As-cast TAC-2

166.7 ± 2.3

57.5 ± 1.8

14.3

As-cast TAC-3

118.0 ± 2.7

63.6 ± 2.4

13.9

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Fig. 6 Compressive engineering stress–strain curves of extruded Mg–6Sn–1Al–xCu (x  0.5, 1, 2) (wt%) alloys Table 4 Tensile mechanical properties of extruded Mg–6Sn–1Al–xCu (x  0.5, 1, 2) (wt%) alloys Alloy

Ultimate tensile strength, UTS (MPa)

0.2% tensile yield strength, TYS (MPa)

Fracture strain (%)

Extruded TAC-1

300.1 ± 3.8

174.8 ± 7.6

7.42

Extruded TAC-2

309.4 ± 4.6

192.3 ± 6.0

6.60

Extruded TAC-3

282.5 ± 3.2

172.5 ± 8.8

7.20

amount of grain boundaries, which will block the movement of dislocations. Therefore, reducing grain size can increase the impediment of dislocation movement so as to improve the strength. However, when adding 2.0 wt% Cu, the mechanical properties of the as-cast alloy decreases. From the microstructure, the grain size of the alloy is found to increase and eutectic structure is formed. In addition, coarse and uneven precipitates distributing at the grain boundary is observed. It cannot play a better role in strengthening due to the separation of the matrix. Therefore, the mechanical properties decrease. The tensile properties of extruded Mg–6Sn–1Al–xCu (x  0.5, 1, 2) alloys are shown in Fig. 5b and Table 4. With the continuous increase of the Cu addition, the tensile properties of extruded alloys increase first and then decrease, which is similar to the variation trend of as-cast alloy’s mechanical properties. The UTS of extruded alloys with 0.5, 1.0, and 2.0 wt% Cu additions are 300.1 ± 3.8, 309.4 ± 4.6, and 282.6 ± 3.2 MPa, and the elongations are 7.42, 6.60, and 7.20%, respectively. The UTS of the extruded Mg–6Sn–1Al–1Cu alloy is the highest, which is 85.6% higher than that of the corresponding as-cast alloy, but the elongation is decreased by 53.8%. When adding 2.0 wt% Cu, the tensile properties of the extruded alloy decreases. Figure 6 reveals the compressive stress–strain curves of extruded Mg–6Sn–1Al–xCu (x  0.5, 1, 2) alloys. Compressive mechanical properties

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Table 5 Compressive mechanical properties of extruded Mg–6Sn–1Al–xCu (x  0.5, 1, 2) (wt%) alloys Alloy

Ultimate compressive strength, UCS (MPa)

0.2% compressive yield strength, CYS (MPa)

Extruded TAC-1

420.1 ± 8.6

Extruded TAC-2

418.1 ± 7.4

79.3 ± 9.8

17.1

Extruded TAC-3

424.2 ± 10.2

110.4 ± 11.8

16.3

72.1 ± 10.4

Fracture strain (%) 16.7

are shown in Table 5. The difference of UCS between extruded alloys is not obvious, with an average of 420.8 MPa, while the compressive yield strength enhances with the increase of the Cu addition, as shown in Table 5. The study shows that although the specimens of extruded alloys exhibit a 45° direction cutoff from the axis after the compression test, compressive stress–strain curves still demonstrate distinct characteristics. The average compressive fracture strain reaches 16.7%, which illustrates that extruded alloys have a higher compression deformation resistance while still maintaining superior ductility to a certain extent. The main reason for the improvement of extruded alloys in mechanical properties is the effect of deformation structure and second phases strengthening. For as-cast alloys, although the addition of Cu can refine grains, there are still some coarse columnar crystals and dendrites in the casting structure, which has adverse effects on the mechanical properties. In addition, defects such as blowholes, porosity, and segregation are unavoidable in casting structure of the alloys. After the extrusion process, these defects are improved significantly and the density of extruded alloy increases. At the same time, the coarse second-phase particles in casting structure are broken to form precipitates with fine dispersion after hot extrusion. The interaction between precipitates and dislocations hampers the movement of dislocations, thereby increasing the plastic deformation resistance. The excessive addition of Cu (ωCu  2.0 wt%) pushes up the amount of precipitates in the extruded structure, but they agglomerated during the hot extrusion process and show a tendency of flow line distribution. This will result in partial stress concentration and unevenness of deformation of the alloys under load, which will eventually lead to a decrease in mechanical properties.

4 Conclusions (1) The addition of Cu can refine the grains of the as-cast alloy and promote the precipitation of the Mg2 Sn phases. However, excessive Cu addition leads to segregation, therefore forming coarse second phases and eutectic structures. (2) Proper addition of Cu can improve the tensile strength of as-cast alloys. The as-cast Mg–6Sn–1Al–1Cu alloy achieves the best comprehensive mechanical properties with the tensile strength of 166.7 ± 2.3 MPa and elongation of 14.3%.

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(3) Hot extrusion process can improve alloys’ casting structure. The microstructure of extruded Mg–6Sn–1Al–xCu (x  0.5, 1, 2) alloy contains uniform and dense recrystallized grains with no distortion. As the amount of Cu addition increases, the recrystallized grains are refined. (4) The mechanical properties of extruded alloys are considerably improved compared to the as-cast ones. The mechanical properties of the as-extruded Mg–6Sn–1Al–1Cu alloy are the best with tensile strength of 300.1 ± 3.8 MPa and elongation of 7.42%. Acknowledgements The financial support for this work is provided by the National Natural Science Foundation of China (Nos. 51571102) and the Shandong Provincial Natural Science Foundation, China (ZR2018LE001).

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Strain Rate Sensitivity of GH4720LI Alloy with Two Initial Microstructures During Hot Deformation Zhi-Peng Wan , Tao Wang, Yu Sun, Lian-Xi Hu, Zhao Li, Peihuan Li and Yong Zhang

Abstract In this study, the effects of initial microstructure on the flow stress, strain rate sensitivity (m), and microstructure during hot deformation of GH4720LI alloy with different initial microstructures were studied using hot compression tests over a wide temperatures range of 1080–1180 °C and strain rates (0.001–10 s−1 ) to a final true strain of 0.8. The results showed that flow stress and deformation mechanisms of the alloys were significantly affected by the γ precipitates. The flow stresses of the two initial microstructures (i.e., microstructures AC and AF) presented typical DRX softening behavior and exhibited nearly a consistent variation trend which was decreased with the increase of temperature. The peak stresses in the microstructure of as-forged samples with smaller initial grain size were lower than microstructure AC when deformation temperature was lower than 1160 °C. While the gap between the two sets of specimens gradually decreased over a temperature of 1160 °C, which was mainly attributed to dissolution of the γ precipitates in alloys. According to the analysis of the strain rate sensitivity values distribution maps with two initial microstructures, the deformation mechanisms of the alloys in various deformation conditions were discussed in detail. Dislocation glide/climb was identified as the dominant deformation mechanism at low temperature, while grain boundary sliding and accommodation was confirmed as the main deformation mechanism at high temperature 1180 °C and low strain rate.

1 Introduction GH4720LI is a high-strength precipitates hardening nickel-based superalloy and has been gaining great attention as discs of gas turbine engines for 750 °C long-term Z.-P. Wan (B) · T. Wang · Z. Li · P. Li · Y. Zhang Science and Technology on Advanced High Temperature Structural Materials Laboratory, AEEC Beijing Institute of Aeronautical Materials, Beijing 100095, China e-mail: [email protected] Z.-P. Wan · Y. Sun · L.-X. Hu National Key Laboratory for Precision Hot Processing of Metals, Harbin Institute of Technology, Harbin 150001, China © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_27

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service [1–3]. The mechanical properties of the alloys are used to affect by deformation rate due to strain rate sensitivity (SRS), and the formability of the alloys can be improved attributed to strain rate hardening [4]. For the alloy in this study, GH4720LI alloy presents much more complicated strain rate sensitive behaviors, corresponding to its high volume fraction and various morphologies of γ precipitates induced complex deformation mechanisms [5]. During hot deformation, the volume fraction of γ precipitates decreases with the increase of temperature [6], and strain rate sensitive behaviors are, therefore, influenced by the temperature. In previous studies, abundant tensile and hot compression tests have been performed to illustrate the effect of strain rate on the deformation behavior, for instance, an asforged Ti–5Al–5Mo–5V–1Cr–1Fe titanium alloy [7], a rolling Ti–22Al–25Nb sheet [8] and a Fe–Mn–C–Si–Al austenitic TWIP steel [9]. In addition, compared to the hot compression parameters, the initial microstructure of alloy plays an important role in affecting the hot deformation behavior of the materials [10, 11]. Hence, a number of beneficial investigations considering the effect of initial microstructure on hot deformation behavior and microstructure evolution have been conducted. Liu et al. [12] investigated the microstructure evolution of coarse grain, fine grain, and mixed grain of U720LI during hot deformation, and the effects of initial microstructure on the deformation and dynamic recrystallization behaviors were also studied. The influence of the morphologies of α phase on flow softening mechanism of Ti–5Al–2Sn–2Zr–4Mo–4Cr during hot compression tests was investigated by Li et al. [13]. Recently, Zuo et al. [10] found that the flow stress was increased with increasing the amount of σ-phase in the initial microstructure for as-cast N08028 Ni-based alloy, and the parameters in the constitutive equations were also dependent on the initial microstructure. According to the investigations mentioned above, it should be noticed that several studies are focused on discussing the effect of the initial microstructure on hot deformation behavior by microstructural observation [14–16]. However, the effects of initial grain size on the flow stress as well as the quantitative description between the value of strain rate sensitivity and the relevant deformation mechanisms of GH4720LI alloy has not been well established, and therefore, it is required to be systematically studied. Also, the morphology of γ precipitates and grain size of the γ matrix significantly affect the deformation mechanisms and the final microstructures [6, 17]. Hence, it is essential to identify the effect of the initial microstructures on the strain rate dependence of the deformation mechanisms in GH4720LI alloy. In this study, the results of investigations on the effects of initial microstructures and deformation conditions on the flow stress and strain rate sensitivity (m) are investigated during the hot compression tests. In addition, microstructure of the alloy, with two initial microstructures, under various processing parameters such as strain rate and temperature is studied in depth.

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2 Experimental The materials used in this study were 410 mm diameter homogenized GH4720LI as-cast ingot and a 120 mm diameter GH4720LI as-forged bar. The measured composition (in wt%) of the as-cast ingot and as-forged bar were Cr 16, Co 14.5, W 1.3, Mo 3, Al 2.5, Ti 5, and balance Ni. The dissolution temperature of γ precipitates was determined to be 1156 °C [12]. Cylindrical samples of 10 mm in diameter and a height of 15 mm applied in hot compression tests were machined from the wrought bar and homogenized as-cast ingot, respectively. Hot compression tests of GH4720LI alloy with two different initial microstructures were conducted on Gleeble 3800 thermo-mechanical simulator over a wide temperature range of 1080–1180 °C and at constant strain rates of 0.01, 0.1, 1, and 10 s−1 . All samples were heated to the deformation temperatures at a heating rate of 10 °C/s, held for 5 min to eliminate the thermal gradient throughout the samples. Then, it is compressed to a strain of 0.8 at a constant strain rate, followed by immediately water cooling. The deformed samples were sliced along the compressed axis direction. The microstructure was examined by using optical microscopy (OM, DM6000M) and scanning electron microscope (SEM, SUPRA55). The samples were chemical etched in a solution of 5 g CuCl2 + 100 ml C2 H5 OH + 100 ml HCl and electro-etched in a solution of H3 PO4 (150 ml) + H2 SO4 (10 ml) + CrO3 (15 g) with a voltage of 5 V and a current of 1 mA. The γ grain size was determined by using the Heyn intercept method [18, 19].

3 Results and Discussion 3.1 Initial Microstructures The initial microstructure shown in Fig. 1a and b are the homogenized GH4720LI as-cast ingot. It can be seen from the figures that microstructure AC consists of a significant number of residue dendrites. Dong et al. [20] suggested that the residue dendrites for the Ni-based alloy can promote the occurrence of the dynamic recrystallization as a result of more recrystallization nucleation sites provided by the residue dendrites regions. Also, the morphology of γ precipitates in as-cast GH4720LI change from spheres over cubes and octocubes to incoherent dendritic structures. Radis et al. [5] considered that the morphology of the γ precipitates depends on the cooling rate, and the several nonspherical, cuboidal and sometimes, even octocube or incoherent dendritic precipitates would form under low cooling rate. The initial microstructure of as-forged GH4720LI alloy is shown in Fig. 1c and d. As shown in the figures that the initial microstructure AF is composed of equiaxed grains and the primary γ precipitates are formed along the initial grains.

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Fig. 1 Micrographs of the initial microstructure of homogenized as-cast ingot (a) (b) and as-forged bar (c) (d) of GH4720LI alloy

3.2 Flow Stress Behavior Flow stress curves determined at different temperatures, strain rates, and initial microstructure are shown in Fig. 2. It is clearly seen from these figures that the peak stress, steady-state stress, and the size of DRX grains are considered to be significantly dependent on the initial microstructures and hot compression test parameters. The flow stresses of the two initial microstructures (i.e., microstructures AC and AF) exhibit nearly a consistent variation trend which is decreased with the increase of temperature. As shown in Fig. 2 that most of the curves present typical DRX softening behavior characterized by a single-peak stress followed by a gradual fall toward a steady-state stress. However, some differences in the flow stress, the shape of flow curve, the peak stress, and the steady-state stress are evident. The flow stress curve of the microstructure AC sample exhibits more obvious peak stress when it deforms at low temperature. For the same hot compression conditions, the flow stress with the initial microstructure of AC is found to be higher than the flow stress of AF. The difference of single-peak flow stress curve associated with the change in the initial microstructure has also been discussed in other investigations [21, 22].

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Fig. 2 True stress–true strain curves at a strain rate of 1 s−1 for various temperatures and initial microstructures of a as-forged bar and b as-cast ingot

It can be seen from Fig. 3 that the difference of initial microstructure has a significant influence on the peak stress of the alloy during hot deformation. The peak stress in the microstructure AF samples with smaller initial grain size is lower than microstructure AC when deformation temperature is lower than 1160 °C, which should be related with the dynamic softening theories. The similar phenomenon for U720LI alloy is also reported by Liu et al. [12]. Compared with the microstructure AC, the volume fraction of γ precipitates in grains for microstructure AF is lower. It has been addressed that the strength of nickel-based superalloy would be enhanced with the increase of γ volume fraction [23]. At lower deformation temperatures, finely γ particles in grains are able to restrict dislocation gliding and rearrangement, leading to an increase of dislocation density and a rise in peak stress. It has been widely acknowledged that DRX is the dominant dynamic softening mechanism for nickel-based alloys during hot deformation. The sample with the microstructure AC is characterized by fine initial grain size which would facilitate the occurrence of DRX behavior according to an increase in the density of nucleation site. Due to the formation of DRX grains and the associated boundary migration of the DRX grain-induced

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Fig. 3 Peak stress of different initial microstructures for various deformation parameters

dynamic softening behavior, a higher decreasing rate of strain hardening can be achieved in the microstructure AC. Also, it is assumed that the dislocation substructures and subgrain boundaries would be consumed by the continuous original boundary migration, leading to a decrease in the dislocation density [6]. Hence, the value of peak stress in the sample with the microstructure AF is less than the AC. The peak stress gap between the two sets of specimens is gradually decreased, when the specimens deform over a temperature of 1160 °C. The thermodynamic equilibrium phase diagram of alloy indicates that the solvus temperature of the main γ phase precipitates is 1156 °C. When the samples are deformed at 1160 °C, the γ precipitates in alloys are dissolved. The pinning effects of γ precipitates on grain boundary migration and dislocation gliding are disappeared [19, 24]. Hence, this difference decreases with the increase of temperature as a result of the dissolution of γ precipitates in alloys.

3.3 Strain Rate Sensitivity The strain rate sensitivity is usually associated with the deformation mechanisms of the material and would significantly influence plastic deformation behaviors, with the range from 0 to 1 [25, 26]. The mechanical properties of the alloys are affected by strain rate sensitivity associated with the strain rate, and the strain rate hardening effects are able to promote the formability of the alloys. Hence, it is of great importance to determine the strain rate sensitivity (m) coefficients for illustrating the dependence of deformation mechanisms on processing parameters during hot deformation. The strain rate sensitivity index (m) can be determined from the following equation [7, 27, 28]:  d lg σ  (1) m d lg ε˙ ε,T

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Fig. 4 Strain rate sensitivity values distribution maps for hot deformation of GH4720LI alloy with initial microstructure of AF (a) and AC (b)

where σ , ε, ε˙ , and T are the flow stress, strain, strain rate, and hot compression temperature, respectively. The strain rate sensitivities of two initial microstructures determined from the true stress–true strain curves of GH4720LI alloy during hot compression tests are presented in Fig. 4. The strain rate sensitivity exhibits a noticeable dependence on deformation temperature, strain rate as well as initial microstructure. It can be seen from Fig. 4a that the value of m for microstructure AC is larger than that of microstructure AF at a strain of 0.8, strain rates of 0.1–1 s−1 and a broad range of temperature encompassing 1120–1180 °C. While the value of m of microstructure AF is found to be larger in a range of strain rates 0.03–1.78 s−1 and low temperatures (i.e., 1080 and 1100 °C). In the temperature range of 1100–1160 °C, the value of microstructure AF is larger or similar to that of microstructure AC at strain rates of

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0.01 and 10 s−1 . This can be reasonably illustrated according to the microstructure evolution. It is well known that the deformation mechanism of dislocation glide and climb occurs at low strain rate sensitivity coefficients conditions (1/6 < m < 1/5), and diffusion accommodated grain boundary sliding or grain boundary diffusion presents at high strain rate sensitivity coefficients conditions (m ≈ 1). The deformation mechanisms of dislocation accommodated GBS or lattice diffusion may occur when the value of m equals to 0.5. Figures 5a and 6a show the microstructures of GH4720LI with different initial microstructures at a deformation strain of 0.8, temperature of 1100 °C and strain rate of 0.1 s−1 . For microstructures AF and AC, m falls into the range of 0.20–0.27 and 0.13–0.17 when deformed in at a temperature of 1100 °C and strain rate of 0.1 s−1 which suggests that the deformation mechanism is likely dislocation glide/climb. As shown in Figs. 5a and 6a that the deformed microstructures are mainly composed of equiaxed and fine γ grains. It is widely acknowledged that the microstructures characterized by fine and equiaxed grains are beneficial for grain boundary sliding and accommodation [29, 30]. While in this study, the pinning effect induced by the primary γ precipitates long boundaries would hinder the

Fig. 5 OM micrographs with the initial microstructure of AC deformed at 1100 °C/0.8 with a strain rate of 0.1 s−1 (a) and 10 s−1 (b) and 1180 °C/0.8 with a strain rate of 0.1 s−1 (c) and 10 s−1 (d)

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Fig. 6 OM micrographs with the initial microstructure of AF deformed at 1100 °C/0.8 with a strain rate of 0.1 s−1 (a) and 10 s−1 (b) and 1180 °C/0.8 with a strain rate of 0.1 s−1 (c) and 10 s−1 (d)

Fig. 7 OM micrographs deformed at 1180 °C/10 s−1 and a strain of 0.8 with the initial microstructure of AC (a) and AF (b)

boundary migration. Hence, the dislocation glide and climb is therefore considered as the dominant deformation mechanism for the alloys deformed at this condition. All of the γ precipitates in materials are dissolved according to the phase diagram of GH4720LI alloy when the deformation temperature increases to 1180 °C,

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which leads to a significant increase of the DRX grain size. The pinning effect on the grain boundaries induced by the γ precipitates disappears. The values of m are larger when deformed at higher deformation temperature of 1180 °C and low strain rate, and the grain boundary sliding is therefore identified as the dominant deformation mechanism for the two initial microstructures. However, the value of m is higher for the microstructure AF. Compared with the microstructure AF, the volume fraction of the grain boundaries in microstructure AC is lower, and nucleation rate of new DRX grains is therefore restrained. This results in the occurrence of localized shearing bands to accommodate the severe plastic deformation in the original individual coarse grain for the microstructure AC. It has been recognized that high deformation temperatures are beneficial for grain boundary sliding and accommodation, while the boundaries are suppressed when the samples deformed at high stain rate. Hence, the values of m for microstructure AF and AC is decreased to 0.13–0.17 at the higher deformation strain rate of 10 s−1 . Often, the grain boundaries function as sites of weakness when the materials deformed at high temperatures [31]. Conceptually, grain boundary sliding may easier occur, which would lead to the opening up voids along the boundaries at a high strain rate to accommodate the incompatible deformation of adjacent grains. Hence, if voids are located preferentially at grain boundaries, the fracture characterized by intergranular fracture modes will occur, as shown in Fig. 7. In addition, the maximum m value of 0.41 presents at a deformation temperature of 1180 °C and strain rate of 0.01 s−1 for microstructure AF, which suggests the occurrence of superplastic deformation behavior [32, 33].

4 Conclusion In this study, the hot compression tests have been performed on GH4720LI alloy with microstructure AF and AC. The flow stresses are significantly dependent on the initial microstructure and hot compression test parameters. The peak stresses in the microstructure AF samples with smaller initial grain size are lower than microstructure AC when deformation temperature is lower than 1160 °C, while the peak stress gap between the two sets of specimens gradually decreases when the specimens deform over a temperature of 1160 °C. Strain rate sensitivity is significantly affected by initial microstructure, deformation temperature, and strain rate, and is associated with the deformation mechanisms of the material. The dominant deformation mechanism is identified as dislocation glide/climb at low temperature. Comparing with the microstructures AF, the occurrence of localized shearing bands to accommodate the severe plastic deformation in the original individual coarse grain for the microstructure AC is observed at high deformation temperature 1180 °C, which will lead to a decrease of m. Acknowledgements This research did not receive any specific grant from funding agencies in the public, commercial, or not-for-profit sectors.

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Microstructure and Properties of Rare-Earth B4 C–Copper Composites Yan Li, Meihui Song, Yu Zhang, Yanchun Li and Xiaochen Zhang

Abstract The B4 C/Cu composites were prepared by powder metallurgy. Scanning electron microscopy, transmission electron microscopy, hardness tester, conductivity meter, and other analytical tests were used to study the effects of alloying elements and rare-earth modification processes on the microstructure and properties of composites. The results showed that the B4 C/Cu composites prepared by powder metallurgy were densely packed, and the Y2 O3 modified layer was uniformly coated on the surface of B4 C particles. The modified B4 C particles had a good interface with Cu and no interfacial reaction occurred. The addition of low melting point alloying element Bi could improve the hardness and electrical conductivity of B4 C/Cu composites, but had little effect on the thermal conductivity. Surface modification of B4 C with rare-earth salt could effectively increase the bonding ability between B4 C particles and Cu, and endowed B4 C/Cu composite materials with excellent physical and mechanical properties. For the composite with 2 wt% B4 C particles, 2 wt% Bi particles, and 96 wt% copper powder, B4 C/Cu composites had good mechanical and physical properties, of which the hardness was 179 HB; the density was 99.6%; the electrical conductivity was 66 %IACS; and the thermal conductivity could reach 262 W/m °C. These properties meet the requirements of electrical contact materials.

1 Introduction With the continuous improvement of manufacturing level of the electric contacts and the increase of the variety, the research on the contact materials such as Cu to take place of Ag has been caused great attention, especially to meet the requirements of saving precious metals such as Ag [1]. In accordance to its conductivity and thermal conductivity, Cu is closest to Ag [2]. However, the main obstacle of Cu to become the electrical contact material is that its surface is easy to oxidize. Moreover, its oxide has very low conductivity, which increases the contact resistance rapidly, and makes the Y. Li (B) · M. Song · Y. Zhang · Y. Li · X. Zhang Institute of Advanced Technology, Heilongjiang Academy of Sciences, Harbin 150020, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_28

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material easy to heat while in use, as well as directly affecting the working reliability of the electrical contact switch [3]. The hardness (HV) of B4 C can reach 3500–4500 [4], which is the third hardest material only behind the diamond and boron nitride in nature. At the same time, it has a constant high-temperature strength [5]. If B4 C is added into Cu as an augmented body to form the B4 Cp/Cu composites, it is expected that the composites will have excellent performance to meet the requirements of electrical contact material. However, B4 C and Cu are not wetted, so as to make it difficult to prepare B4 Cp/Cu, and usually, it requires a higher preparation temperature [6, 7]. Even so, its uniformity of composites is poor and its performance is far below the theoretical value. The previous studies have shown that rare-earth elements have a special chemical structure. The introduction of a small amount of rare-earth oxides in the composite can improve the wettability of ceramic particles and aluminum alloys, so that the metal matrix composites have excellent physical and mechanical properties [8]. Besides, the introduction of rare-earth elements in sintering of ceramic materials is a commonly used method to reduce sintering temperature and improve material properties [9]. However, this method is rarely used in the study of B4 Cp/Cu composites. Through introducing rare-earth elements or rare-earth compounds into B4 Cp/Cu composites, the wettability and interfacial bonding strength of B4 C–Cu interface will be improved and the properties of the composites will be improved, as well. However, none has been reported so far [10]. According to the requirements of the new type of electrical contact material, this paper is based on the situations, such as easy oxidation, low mechanical properties, poor welding performance in the atmosphere, and so on [11]. By using the composite strengthening principle, this paper optimized the design of composite components and preparation process while taking metal copper as the matrix. Mechanical properties, resistance to welding, and oxidation resistance can be improved by adding components of different physical properties [12]. At the same time, a new type of B4 Cp/Cu composite with rare-earth elements was developed by high-energy ball milling and powder metallurgy sintering during the process of preparation [13]. Moreover, the microstructure and interface state of the composites were studied by means of various analytical methods, as well as the effects of B4 C with modification of rare-earth salt on the properties of the composites were investigated.

2 Experimental Materials and Experimental Methods 2.1 Preparation of the Modified B4 C Y(NO3 )3 ·6H2 O solution (1 mol/L) was configured with anhydrous ethanol as a solvent, and then B4 C powder was added in a certain proportion (according to the ratio of the quality of B4 C and the rare-earth solution was 1 g:10 ml). After mixing and drying, it was sintered at 500 °C for 1 hour in the muffle furnace, and the B4 C powder with modification of rare-earth salt was obtained.

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2.2 Composite Preparation Composite materials were prepared by powder metallurgy. In a certain proportion, the modified B4 C powder and Bi powder were added to the electrolytic copper powder of 50 µm. Through the exploration of the previous experiments, the best ratio was 2% for Bi powder and 2% for modified B4 C. Then, the powders were mixed under the protection of nitrogen. After the pressing of 500 MPa, the B4 Cp/Bi/Cu composite was obtained after sintering under 900 °C in vacuum.

2.3 Test and Characterization The hardness was measured by HBRV-187.5 type of Brinell hardness gauge. The conductivity of composite material was measured by SIGMATEST 2.069 eddy current conductance instrument, and the test frequency was 60 kHz. The conductivity unit of the test material was %IACS. On the S4700 scanning electron microscope (SEM) and the JEM-2100 transmission electron microscope (TEM), the microstructure of the composite was observed, and the accelerated voltage of the transmission electron microscope was 200 kV and the length of the camera was 300 mm.

3 Microstructure 3.1 Morphology of the Original Material Figure 1 is the photograph of the original morphology of (a) Bi powder; (b) B4 C powder; (c) B4 C powder with modification of yttrium nitrate; and (d) copper powder. In the figure, the copper powder used in this paper is a dendritic structure prepared by the electrolysis process. The modified boron carbide powder surface formed a coating. The coating was reacted and sintered through yttrium nitrate–ethanol solution. After that, the Y2 O3 coating is formed on the surface of boron carbide.

3.2 SEM of Composite Materials The ball milling process could reduce the particle size, and change the shape of the particles, as well as mixing the components fully [14]. However, the composite material contained metal Bi, which was very easy to oxidize in the air, so the ball milling time could not be too long. In this paper, the ball milling time is 2 hours. Figure 2 was the image of the powder after ball-milling.

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Fig. 1 SEM of original powders a Bi, b B4 C, c modification B4 C, and d copper Fig. 2 SEM of the mixed powder

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Figure 3a is B4 C/Bi/Cu without modification of yttrium nitrate. Figure 3b is B4 C/Bi/Cu with modification of yttrium nitrate. It can be seen from the diagram that there were many gaps in the unmodified material, which resulted in a decrease in the comprehensive ability of the material. On the other hand, the gap of the modified B4 C/Cu material decreased obviously. This is because the non-wetting and bonding ability of the B4 C powder and Cu powder was not strong. The rare-earth elements have a special chemical structure. The use of rare-earth salt to modify the B4 C powder could effectively enhance the binding ability of the B4 C powder with Cu, so as to reduce the gaps in the material and improve the properties of the material. Figure 3c is the energy spectrum of the modified material, which further proved the existence of yttrium oxide in the material.

Fig. 3 SEM of composite materials a B4 C/Bi/Cu, b B4 Cp/Bi/Cu, and c energy spectrum analysis of B4 Cp/Bi/Cu

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4 TEM of Composite Materials 4.1 Interface of Composite Materials The interface is an important part of composite materials, and its properties determine the properties of materials directly [15]. Figure 4 are the interfaces of B4 C/Cu composite, B4 C/Bi/Cu composite, and B4 Cp/Bi/Cu composite. It can be seen from Fig. 4b and c that the interface of the composite had a good interface and no cracking, while the B4 C and Cu interfaces in (a) had obvious cracking phenomena. It indicates that the binding was not ideal.

Fig. 4 Interface of composite materials a B4 C/Cu, b B4 C/Bi/Cu, and c B4 Cp/Bi/Cu

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4.2 Crystal Defect Figure 5 is a photo of the crystal defect of the composite. It can be seen from the diagram that there were a large number of dislocations and twins in Cu. The main dislocations were the edge dislocations. This was because the pressing pressure

Fig. 5 Dislocations and twins in B4 Cp/Bi/Cu composite materials

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and sintering heat could cause plastic deformation of Cu during the fabrication of composite materials. Therefore, it resulted in a large number of edge dislocation and deformation twins. Defects such as dislocations and twins in composites were beneficial to improve the mechanical properties of materials, but they would also adversely affect the thermal conductivity and conductivity of materials. Besides, the addition of B4 C enhanced phase could improve the hardness and wear resistance of the material, mainly due to the dispersion and dislocation strengthening of B4 C.

5 Performance Analysis Table 1 shows the basic properties of unmodified B4 C/Cu and modified B4 C/Cu composites. As stated in the table, the addition of Bi powder to the composite affected the thermal conductivity, electrical conductivity, and hardness of the material. This is because Bi could be completely melted in the copper matrix during the preparation process. After cooling, Bi would be distributed uniformly in the matrix. The structure of the material would be formed into “steady state” when the material was covered by the surface of the material. The “steady state” means that the hardness and electrical conductivity of the material were improved. B4 C itself has high hardness and can improve the oxidation resistance of the matrix [16]. Therefore, B4 C was chosen to be an enhanced phase. However, the wettability of copper and B4 C is not good. According to the previous study, the rareearth element had a special chemical structure. The introduction of a small amount of rare-earth oxides in the composite could improve the wettability of ceramic particles and copper alloys, so that the metal matrix composites have excellent physical and mechanical properties. The method had been applied in the research of AlNp/Cu composites.

Table 1 Properties of composite materials Materials

Density (kg/m3 )

Relative density (%)

Thermal conductivity (W/m °C)

Conductivity Hardness (%IACS) (HB)

B4 C/Cu

7.34

98.3

118

46

149

B4 C/Bi/Cu

7.47

98.6

163

54

138

B4 Cp/Bi/Cu

7.95

99.6

262

66

179

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6 Conclusion (1) When yttrium nitrate is used to modify the surface of B4 C, Y2 O3 -modified layer is formed on the surface, which can effectively improve the wettability of Cu. (2) B4 Cp/Bi/Cu composites are prepared by powder metallurgy. The compactness of the composites reaches 99.6%, and the hardness is up to 179 HB. Meanwhile, the conductivity is 66 %IACS, and the thermal conductivity can reach 262 W/m °C. (3) The interface of Cu–B4 C is well bonded without cracking and interfacial reaction, which is well bonded with Bi particles. (4) There are a large number of edge dislocations and twins in the composites, which enhance the mechanical properties but reduce the electrical properties of the composites.

References 1. H.-R. Zhao, W. Luo, Y.-Z. Xie et al., Study advance of contact material in 2015. Electr. Eng. Mater. 1, 24–28 (2016) 2. R. Jamaati, M.R. Toroghinejad, Application of ARB process for manufacturing high-strength, finely dispersed and highly uniform Cu/Al2 O3 composite. Mater. Sci. Eng. A 527(27), 7430–7435 (2010) 3. Y. Jiang, M.-G. Yang, Z.-C. Yang et al., Research status of copper-based electrical contact materials for medium and high voltage apparatus. Electr. Eng. Mater. 4, 27–30, 35 (2013) 4. S.G. Savio, K. Ramanjaneyulu, V. Madhu, T. Balakrishna Bhat, An experimental study on ballistic performance of boron carbide tiles. Int. J. Impact Eng. 38, 535–541 (2011) 5. K. Rajkumar, K. Kundu, S. Aravindan et al., Accelerated wear testing for evaluating the life characteristics of copper-graphite tribological composite. Mater. Des. 32, 3029–3035 (2011) 6. J. Zhang, L. He, Y. Zhou, Highly conductive and strengthened copper matrix composite reinforced by Zr2 Al3 C4 particulates. Scripta Mater. 60, 976–979 (2009) 7. X. Tang, Hot Deformation Behavior and Hot Extrusion Process of ZnSnO3 /Cu Electrical Contact Materials (Harbin Institute of Technology, 2015) 8. Z.Y. Xiao, M.Y. Ke, L. Fang et al., Die wall lubricated warm compacting and sintering behaviors of pre-mixed Fe–Ni–Cu–Mo–C powders. J. Mater. Process. Technol. 209, 4527–4530 (2009) 9. G.-L. Li, X.-C. Jiang, M. Wen et al., Studies of B4 C particles reinforced copper matrix composite. J. Mater. Eng. 8, 32–35 (2001) 10. J.-P. Li, S.-H. Meng, J.-C. Han, Structure and flaws of CuCr alloys by explosive compaction. J. Harbin Inst. Technol. 12(2), 135–138 (2005) 11. S.J. Sun, S. Sakai, H.G. Suzuki, TEM observation of Cr fibers in Cu–15Cr–0.5Fe in situ composites. Mater. Trans. 41(5), 613–616 (2000) 12. Z. Mu, H.-R. Geng, M.-M. Li et al., Effects of Y2 O3 on the property of copper based contact materials. Compos. B 52, 51–55 (2013) 13. C.J. Tu, D. Chen, Z.H. Chen, Improving the tribological behavior of graphite/Cu matrix selflubricating composite contact strip by electroplating Zn on graphite. Tribol. Lett. 31, 91–98 (2008)

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14. Y. Li, M. Song, Y. Zhang et al., Influence of the ball mill process on titanium alloy powder particle size. Chem. Eng. 06, 13–15, 53 (2017) 15. Y. Li, M. Song, Q. Yu, X. Zhang, Improvement of anti-hydrolytic property of AlN powder modified by the rare-earth salts. Mater. Sci. Forum 816, 9–14 (2015) 16. M.K. Habibi, A.S. Hamouda, M. Gupta, Hybridizing boron carbide particles with aluminum to enhance the mechanical response of magnesium based nano-composites. J. Alloy. Compd. 550, 83–93 (2013)

Evaluation of Microstructure and Refining Effect of Al–TiB2 and Al–5Ti–1B Grain Refiners Mengke He, Lihua Chai, Hongda Wang, Ziyong Chen and Yapeng Cui

Abstract High-concentration Al–Ti–B master alloy was prepared by melt selfpropagating high-temperature synthesis, and then the extruder bar of intermediate alloy was also prepared. The grain refiner can be divided into two categories. One is TiB2 particles only, and the other is two particulate matter containing both TiB2 and TiAl3 . The microstructures of the two-grain refiners were observed under scanning electron microscope (SEM) and compared with the commercial refiners. The fining effects of the two-grain refiners were analyzed. The refining mechanism of TiAl3 and TiB2 was discussed, and the optimum dosage and the refinement scheme of pure aluminum were optimized.

The addition of grain refiners in aluminum melt can effectively refine the grain, thus improving the strength and plasticity of the material. At the present stage, the Al–Ti–B grain refiner still has some problems in different degrees, such as composition segregation, uneven microstructure, and unstable performance [1]. Besides, the purity and the stability of the Al–Ti–B grain refiner need to be improved [2]. At the same time, Al–Ti–B grain refiners contain TiAl3 structure, the effect of which on the refinement of aluminum alloys is not yet possible to determine [3, 4]. Scholars from all over the world have done a lot of researches on the refining mechanism of Al–Ti–B alloy. Because of the opaque metal, it is impossible to observe the nucleation process of the grain directly [5]. Due to the size of the nucleus being submicron, it is difficult to distinguish the intermetallic compounds in the alloy, and the nucleation element may react with the matrix element during the nucleation process [6]. The refining mechanism of grain refiners mainly includes paritectic phase diagram theory, particle theory, the value-added theory of α-Al and the theory of double nucleation [7, 8]. At present, Al–5Ti–1B refiner is the main grain refiner in the world aluminum industry. Because of the effective nucleation of TiB2 particles, some scholars use high content Al–TiB2 master alloy to refine aluminum alloy, and it is proved that Al–TiB2 also has a good refining effect on pure aluminum [8]. However, the refining effect and M. He · L. Chai · H. Wang · Z. Chen (B) · Y. Cui College of Materials Science and Engineering, Beijing University of Technology, Beijing 100124, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_29

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refining mechanism of the two kinds of refiners were not evaluated and analyzed. Therefore, it is necessary to determine the refinement effect of Al–TiB2 refiner and compare it with Al–5Ti–1B refiner. In this paper, X-ray diffractometer (XRD), metallographic analysis (OM), scanning electron microscopy (SEM), and energy spectrum analysis (EDS) were used to analyze the monopoly of the grain refiners, aiming at comparing the refining effect of Al–5Ti–1B grain refiner and Al–TiB2 grain refiner [9]. With low-cost titanium dioxide, boric acid, and aluminum powder as raw materials, Al–5Ti–1B grain refiner and Al–TiB2 grain refiner were prepared by direct melt self-propagating synthesis method. The microstructure and refining effect of the two kinds of grain refiners were compared and the refining mechanism of TiAl3 phase and TiB2 phase was explored.

1 Experimental Materials and Methods A graphite crucible resistance furnace was used to heat up and the grain refiner was prepared by melt self-propagating reaction. The raw materials used are: pure aluminum–aluminum ingot, titanium dioxide powder, boric acid powder, aluminum powder, titanium powder, and hexichloroethane refining agent. Among them, the mole ratio of the selected Al–TiB2 refiner ingredients is: Ti/B  1:4, Ti/TiO2  2:3, Al–5Ti–1B, the molar ratio of the refiner is Ti/B  1:2, Ti/TiO2  2:3. After mixing aluminum powder, titanium powder, boric acid powder, and titanium dioxide powder in proportion, they are mixed uniformly by machine and pressed into cylindrical powder with a diameter of 62.5 mm and height of 20 mm. After melting pure aluminum ingot at 900 degree, the pre-prepared powder was pressed into the graphite bell jar, using mechanical stirring for 20 min. After the reaction ended, C2 Cl6 degassing refining was pressed into the melt to remove the surface dross, and the refiner ingot was obtained in a column steel die with a diameter of 90 mm and a height of 300 mm at 800° (the steel die was warmed at 200 °C for 1 h before). The high-temperature extrusion of grain refiner ingots was carried out. The extrusion temperature was 430 °C, the extrusion ratio was 89.7:1, and the extrusion rate was 5–8 mm/s. The microstructure of extruded grain refiners was analyzed, and the size, morphology, and distribution of the phase were observed. According to previous studies, the fining agent can achieve the best refining effect at 2–5 min. On the basis of this work, the best contact time is chosen to refine the extruding rod grain refiner. The contact time is 2 min. The samples for the refined experiment were polished to the mirror with 200–2000# sandpaper. The surface corrosion (corrosion time was 3–5 s) was carried out in 60%HCL + 30%HNO3 + 5%HF + 5H2 O (volume fraction), and the macro photo was taken after corrosion. The grain size of the refined aluminum was measured according to the GB/T63942002 standard. The average size of grain is determined by cutoff point method. In order to obtain a reasonable average value, 3–5 representative field of view is selected

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to measure the number of grain cut by a straight line in each market, and the average intercept D is used to determine the grain size. D  L/(M · P)

(1)

In which L is the measured length of the line segment, mm, M for observation magnification, and P is the number of cutoff points on the measured line segment. The D8 ADVANCE X-ray diffractometer (XRD) was used to analyze the phase of the refiner, and the FEI QUANTA FEG 650 scanning electron microscope (SEM) and the backscattered electron image were used to observe the ion morphology and distribution of the refiners. The elemental composition of the microelements was qualitatively and semiquantitatively analyzed by means of X-ray energy dispersive spectrometer (EDS) equipped with SEM.

2 Results and Analysis 2.1 Phase Analysis of Al–5Ti–1B Grain Refiners and Al–TiB2 Grain Refiners The XRD analysis of Al–5Ti–1B refiner ingot was carried out. As shown in Fig. 1, the Al–5Ti–1B is mainly containing α-Al, TiB2 and TiAl3 phases. Besides, the AlB2 phase cannot be found in the spectrum, which indicates that the maximum of Ti and B elements transfer to TiAl3 and TiB2 phase, and the quantity of the Al–5Ti–1B is sufficient. The experimental scheme is feasible, that is, the melt self-propagating method is adopted at 950 °C, and the method of replacing the partial oxide with a single substance is effective and feasible. XRD analysis of Al–TiB2 refiner ingot is carried out. As shown in Fig. 2, the sample has only TiB2 phase except for the alphaAl phase, and the sample does not appear AlB2 and TiAl3 phase, which indicates that the maximum of Ti and B elements is transformed into the TiB2 effective phase, and the quantity is sufficient.

2.2 Microstructure Analysis of Al–5Ti–1B Refiner and Al–TiB2 Refiner 2.2.1

Microstructural Analysis of Al–5Ti–1B Grain Refiner

Scanning electron microscopy was used to observe the microstructure of the samples, and the distribution of TiB2 and TiAl3 and particle size were also studied. Figure 3 is SEM scanning micrographs of new Al–5Ti–1B grain refiner sample.

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Fig. 1 XRD spectrum of Al–5Ti–1B grain refiner ingot

Fig. 2 XRD spectrum of Al–TiB2 grain refiner ingot

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Fig. 3 Scanning microstructure of Al–5Ti–1B grain refiner

From Fig. 3a, b, one can see that there are two kinds of second phases distributed on the matrix, which are bar-shaped phase and particles. The bar-shaped phase is uniformly dispersed, however, the particle phase distributes at grain boundaries and some of which aggregates as cluster in local areas as seen in Fig. 3c. Figure 3d is a morphographic map of the particle phase at a larger magnification. The cluster is more concentrated, and the shape and size are more uniform. It is known from the literature [10] that TiB2 phase is formed at the intergranular and grain boundaries, which can obstruct the dendritic growth of the matrix and thus effectively refine the cast microstructure. Hyman and others through a series of studies on Al–5Ti–1B alloys showed that when the content of B is less than 1%, the shape of the TiAl3 is thick strip, needle-like and thin skin, and when the content of B is greater than 1%, the morphology of the TiB2 is massive and granular. TiB2 in Fig. 3c is granular, which indicates that the B content in the refiner sample is more than 1%, and the burning of B is less. In general, the second phase distribution is relatively uniform, the number is enough, and the local aggregation state occurs. Energy spectrum analysis Fig. 4a shows that there are two kinds of particles with different morphologies in the refiner, “A,” “B,” and “C” three points are selected for

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(a)

(b) C

B

A

(c)

(d)

Fig. 4 EDS analysis of precipitations in Al–5Ti–1B refiner

analysis, respectively. “A” points are lump particles, the size of which is below 2 mm. The energy spectrum scanning analysis (Fig. 4b) shows that the particle is mainly composed of Ti, B, and trace Al. The percentage of Ti atom and B atom is close to 1:2, so the granular particles are determined to be TiB2 phase combined with XRD analysis. The “B” point is a part of the rod-like particles. From the analysis results, as seen in Fig. 4c, two elements of Ti and Al can be seen, in which the percentage of the Ti atom and the Al atom is close to 1:3, thus the bar-like phase is further determined to be the TiAl3 phase. The “C” point is black massive microstructure, which is mainly composed of two elements of Al and O as the energy spectrum analysis shown in Fig. 4d. The percentage of Al atom and O atom is close to 2:3. It is further speculated that the black bulk phase is Al2 O3 phase, which is considered as an oxide impurity. The energy spectrum analysis in Table 1 shows that the Ti element content in the refiner is about 5 wt% and the content of the B element is about 1 wt%, and the refiner can be used as Al–5Ti–1B refiner with the combination of Fig. 3.

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Table 1 ICP-MS analysis of the matrix of Al–5Ti–1B refiner Sample number

Original sample number

B (w/%)

Ti (w/%)

001

1

0.82

4.96

002

2

0.86

5.03

003

3

0.83

5.01

Fig. 5 Scanning microstructure of Al–TiB2 grain refiner

2.2.2

Microstructural Analysis of Al–TiB2 Grain Refiner

Figure 5 shows SEM scanning micrographs of the Al–TiB2 grain refiner sample. Only one kind of particles was found in the matrix. The morphology of the second phase particles can be seen in Fig. 5a and b. The number of TiB2 phases is relatively high and the distribution is relatively uniform. In Fig. 5c, TiB2 is granular, which indicates that the B content in the Al–TiB2 refiner sample is more than 1%. Figure 5d shows the morph graphic images of the TiB2 phase at larger magnification, in which The TiB2 phase of the cluster is more dispersed, with a certain distance, the particle size is also uniform. In general, the second phase

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Table 2 ICP-MS analysis of Al–TiB2 grain refiner Sample number

Original sample number

B (w/%)

001

1

4.89

Ti (w/%) 9.97

002

2

5.11

10.32

003

3

4.96

10.06

Fig. 6 Microstructure of pure aluminum ingots with different addition of Al–5Ti–1B refiners. a Pure aluminum without adding refiner, b adding 0.1 wt% refiner, c adding 0.15 wt% refiner, d adding 0.2 wt% refiner, e adding 0.25 wt% refiner

distribution is more homogeneous, and the quantity is more than the TiB2 content of Al–5Ti–1B grain refiner. The energy spectrum analysis of the Al–TiB2 refiner was carried out in a large field of vision. As shown in Table 2, the content of Ti element in the refiner is about 10 wt% and the content of B element is about 5 wt%, and the refiner could be used as Al-15 wt%TiB2 refiner with the combination of Fig. 6.

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Fig. 7 Grain size map of pure aluminum refined with different Al–5Ti–1B refiners and Al–TiB2 refiners

3 Refining Effect and Refining Mechanism of Al–5Ti–1B Grain Refiner and Al–TiB2 Grain Refiner 3.1 Effect of Addition Amount on Refining Effect of Al–5Ti–1B Grain Refiner As can be seen from Fig. 6, adding a small amount of Al–5Ti–1B refiner in pure aluminum can get a clear refining effect. According to the GB/T6394-2002 standard, the size of the aluminum grain refined is measured. The average grain size was measured by the intercept method as the basis for evaluation, and the results are shown in Fig. 7. When the grain refiner addition amount is 0.1 wt%, the grain size decreases to 2 mm immediately, but the grain size of the refined aluminum is not uniform. When the refiner added is 0.15 wt%, the grain size could reach 1.01 mm, which decreases by 76.3% compared with the grain size of pure Al, and the grain size of the refined pure aluminum tended to be uniform. After adding Al–5Ti–1B refiner is up to 0.2 and 0.25 wt%, the refined grain size of pure aluminum is 0.98 and 0.99 mm, respectively, which has not been much changed as compared with the addition amount of 0.15 wt%. Therefore, the optimum addition of Al–5Ti–1B refiner is 0.15 wt%, and the refinement effect of the refiner will not be greatly improved after adding more refiners. From Fig. 6, it can be seen that some of the refined aluminum ingots are not uniform in grain size, grains in the middle are refined but others at margin are coarse. The reason for this phenomenon may be that there is no fully stirring in refining, and the distribution of refiners is uneven, which leading to that the grain size distribution is different. However, the refinement effect and potential of refiners can still be deduced.

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Fig. 8 Microstructure of pure aluminum ingots with different addition of Al–TiB2 refiners. a Pure aluminum without adding refiner, b adding 0.1 wt% refiner, c adding 0.15 wt% refiner, d adding 0.2 wt% refiner, e adding 0.25 wt% refiner

3.2 Effect of Addition Amount on Refining Effect of Al–TiB2 Grain Refiner Figure 8 shows the microstructure of pure Al after adding different Al–TiB2 refiners, clear refining effect of pure aluminum is observed in the alloy with a small amount of Al–TiB2 refiner. The relationship between average size and content of Al–TiB2 refiner exhibits in Fig. 8. When the addition amount is 0.1 wt%, the grain size decreases to 0.96 mm immediately, and the grain is refined obviously. When the addition amount is 0.15 wt%, the grain size is reduced to 0.78 mm, and the grain size is reduced by 77.5% compared with that without refining. After adding the Al–TiB2 refiner up to 0.2 and 0.25 wt%, the refined grain size of pure aluminum is 0.82 and 0.79 mm, respectively, which is not much changed compared with the addition amount of 0.15 wt%. Therefore, the optimum addition of Al–TiB2 refiner is 0.15 wt%, and the refinement effect will not be greatly raised as the addition of the Al–TiB2 refiner higher than that content. From the microstructure shown in Fig. 8, it can be seen that there is a certain grain size inhomogeneous in pure aluminum ingot after refining, owing to the distribution of the refiner is uneven, results in the difference in grain size distribution, but the refining effect and potential of the refiner can be deduced. Compared with Al–5Ti–1B refiner, Al–TiB2 refiner can be added in the same amount to the finer tissue, and the thinning effect of the Al–TiB2 refiner is better under the same amount of addition.

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The refining process of Al–5Ti–1B is as follows: adding Al–5Ti–1B master alloy into Al melt, TiB2 has a high melting point of its own, which makes the particle thermodynamically stable and almost insoluble in Al melt because of its strong B-B bonding force, while TiAl3 has a small bonding force of its strongest Ti–Al bond, and Ti contains Ti in Al melt. The amount is very low (about 0.005–0.01%), so its thermodynamic instability will be partially dissolved. With the prolongation of residence time in aluminum melt and the increase of temperature of Al melt, the dissolution trend of Al melt becomes larger. Then, during solidification, a large number of TiB2 coated with TiAl3 thin layer will be used as the effective nucleation substrate of a-Al, promoting nucleation. At the same time, peritectic reaction occurs in the residual TiAl3 phase in the aluminum melt, which also plays a direct role in nucleation. Therefore, Ti/B > 2.2 is the prerequisite for obtaining the nucleation efficiency of TiB2 . If Ti/B is less than 2.2, no dissolved Ti will form TiAl3 on the surface of TiB2 in Al melt, so the nucleation rate of TiB2 cannot be improved. Second, the second phase in Al–5Ti–1B should be fine and dispersive. The finer TiAl3 and TiB2 particles in the master alloy are, the more TiB2 particles with TiAl3 thin layer are provided in the refining process. At the same time, the more TiAl3 particles remain, and the more nucleated substrate is, so the refining is better. In summary, grain refinement is a complex process, TiAl3 and TiB2 phases have different roles, Ti element can refine grain, B element can enhance the refinement of Ti, TiB2 alone also has refinement effect, and is compared with TiAl3 and TiB2 refiners in the joint effect of weak decay in the holding time, there is still very good refining effect. The holding time in industrial production will be much longer than that in experiments, so Al–TiB2 refiner may have a greater industrial prospect.

4 Conclusions The microstructure and refining effect of two kinds of refiner of Al–5Ti–1B and Al–TiB2 are employed in this paper. (1) There are two kinds of precipitations in the microstructure of Al–5Ti–1B refiner, including rod and needle-like TiAl3 and granular TiB2 , of which the content of TiB2 particles is higher, and the whole distribution is uniform without the obvious state of aggregation. There is only TiB2 phase precipitation in the microstructure of Al–TiB2 refiner, and there is a certain phenomenon of segregation. (2) Refining pure aluminum by using the Al–5Ti–1B refiner rod and the Al–TiB2 refiner rod has obvious refining effect. When the refining time is 2 min, the Al–5Ti–1B refiner rod can refine the pure aluminum grain to 1.01 mm. When the addition amount is 0.15 wt%, the grain size will not change too much with the increase of the additional amount. Al–TiB2 refiner rod can refine the grain of pure aluminum to 0.96 mm when the contact time is 2 min with the addition of 0.1 wt%, and the refining effect is more obvious with increasing of adding

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amount. The grain of pure aluminum can be refined to 0.79 mm when the amount of addition is 0.25 wt%.

References 1. H.B. Michael Rajan, S. Ramabalan, I. Dinaharan, S.J. Vijay, Synthesis and characterization of in situ formed titanium diboride particulate reinforced AA7075 aluminum alloy cast composites. Mater. Des. 44, 438–445 (2013) 2. Y. Ma, Z. Chen, M. Wang, D. Chen, N. Ma, H. Wang, High cycle fatigue behavior of the in-situ TiB2 /7050 composite. Mater. Sci. Eng., A 640, 350–356 (2015) 3. C. Dong, L. Yongkang, B. Liang, M. Naiheng, L. Xian-feng, W. Haowei, Mechanical properties and microstructure of in situ TiB2 7055 composites. Chin. J. Aeronaut. 19, S66–S70 (2006) 4. Zhong L-h, Zhao Y-t, Zhang S-l, G. Chen, S. Chen, Liu Y-h, Microstructure and mechanical properties of in situ TiB2 /7055 composites synthesized by direct magnetochemistry melt reaction. Trans. Nonferrous Metals Soc. China 23(9), 2502–2508 (2013) 5. D. Chen, C. Zou, Y.J. Zhang, N.H. Ma, H.W. Wang, Tensile properties of 15wt. %TiB2 /7055 composite fabricated by in situ method. Adv. Mater. Res. 842, 165–169 (2013) 6. C. Yuanshenru, S. Jienlin, Ageing behaviour of SiCp-reinforced AA 7075 composites. J. Mater. Sci. 32, 1741–1747 (1997) 7. N.V.R. Kumar, E.S. Dwarakadasa, Effect of matrix strength on the mechanical properties of Al–Zn–Mg/SiC P, composites. Compos. A Appl. Sci. Manuf. 31(10), 1139–1145 (2000) 8. N.V.R. Kumar, E.S. Dwarakadasa, Effect of matrix strength on the mechanical properties of Al–Zn–Mg/SiCP composites. Compos. A Appl. Sci. Manuf. 31, 1139–1145 (2000) 9. W. Yuan, J. Zhang, C. Zhang, Z. Chen, Processing of ultra-high strength SiCp/Al–Zn–Mg–Cu composites. J. Mater. Process. Technol. 209, 3251–3255 (2009) 10. Das DK, Mishra PC, Chaubey AK, Singh S. Fabrication process optimization for improved mechanical properties of Al 7075/SiCp metal matrix composites. Manag. Sci. Lett., 297–308 (2016)

Effects of Graphene Content and Aging Process on Mechanical Properties and Corrosion Performance of an A356.2 Aluminum Matrix Composite Kang Wang, Jinfeng Leng, Ran Wang and Shaochen Zhang

Abstract A356.2 aluminum alloy is widely used in the aerospace and automotive industries due to its superior casting ability, high specific strength, and corrosion resistance. Graphene is used in aluminum-based composites for its excellent mechanical properties and unique two-dimensional structure. In this paper, graphene and aluminum powder were mixed uniformly and then graphene-reinforced A356.2 aluminum matrix composites were prepared by atmospheric casting. Research results showed that the hardness change in the age hardening process tended to be faster and then slower. With the increase of graphene content, the peak-aging time of Gr/A356.2 aluminum matrix composites was shortened and the hardness value increased. The peak-hardness was 126HB when aged at 180 °C for 3 h. The addition of graphene improved the pitting resistance of A356.2 aluminum matrix composites in 3.5 wt% NaCl solution, mainly showing the positive shift of pitting potential and the decrease of corrosion current density. The composites under different aging conditions exhibited similar polarization characteristics, indicating the heat treatment not change the electrochemical corrosion behavior of the composites, while the pitting corrosion resistance under the underage condition was the best. The pitting potential was − 0.735 V and the passivation region was the longest at 0.617 V. This was mainly because the second-phase precipitates increased with the aging time, subsequently, more interfaces between the second-phase precipitates and the aluminum alloy supplied the prior sites for pitting.

1 Introduction Aluminum matrix composites have strong vitality materials that have emerged in response to the needs of modern scientific development. Due to excellent casting properties, good corrosion resistance, high specific strength, and low manufacturing cost, Cast aluminum alloy are increasingly used in the automotive and aerospace K. Wang · J. Leng (B) · R. Wang · S. Zhang School of Materials Science and Engineering, University of Jinan, No. 336, West Road of Nan Xinzhuang, Jinan 250022, People’s Republic of China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_30

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industries [1, 2]. In order to reduce environmental pollution and resources crisis, auto lightweight becomes very important. Al–Si alloys have been widely adopted because they reduce product weight, display excellent cast ability and have high specific strength [3]. A356 aluminum alloy is developed in the 1860s as a new type of aluminum alloy and began to use in automobiles hub in the 1970s, which makes the widely application of A356 aluminum alloy [4]. In the last few years, especially particulate-reinforced aluminum matrix composites, are highly regarded because of their remarkable physical and mechanical properties. But they suffer from inadequate ductility [5, 6], uniformity of distribution, and interface reaction. It is a well-known bottleneck that limits the widespread engineering application of micro-composites. To attain a higher strength and retain character of the composite, nano-sized particles are gradually used [7, 8]. Graphene, a one-atom-thick, two-dimensional sheet of carbon atoms is attracting the attention of researchers. Graphene is one of the materials with the highest known strength, it has good toughness and can be bent. The theoretical Young’s modulus of graphene reaches 1.0 TPa and the inherent tensile strength is 130 GPa [9–12]. Therefore, graphene is an effective reinforcement for metal matrix composites. Compared with traditional hard particles materials, graphene is hard and flexible. In addition, it is also characterized by atomic thickness which is smaller than most precipitates in the alloy [13, 14]. There are many studies in the literature on the mechanical properties of graphene-reinforced polymers, but there is scarcity of studies on graphene-reinforced metal nanocomposite. This is likely a result of the greater difficulties in the dispersion and fabrication, and the unknown interfacial chemical reactions in metal composites. Recently, Wang et al. [15] fabricated Al composites reinforced with graphene based on flake powder metallurgy, by adding 0.3 wt% of graphene in Al matrix the ultimate strength of resulting composite enhances by 62%, but there is scarcity of studies in other fields. Due to their wide applications, they come in contact with acids or bases frequently during pickling, de-scaling, electrochemical etching and extensively used in many chemical process industries. Most of the reported studies were conducted on corrosion of various metals and alloys in HCl and H2 SO4 media. In this paper, we mainly investigate the corrosion behavior of different graphene additions and different aging times in 3.5% NaCl solution. It is of great significance to study the corrosion performance on the service life, failure time and protection of composite materials.

2 Experimental Details Al, Mg, Gr, Al–20%Si master alloy were used to produce graphene-reinforced A356 aluminum matrix composites. The ratio of each component is shown in Table 1. First, pure aluminum and Al–20%Si master alloy were put into resistance furnace at 400 °C and melted at 700 °C. The ingot casting would melt completely at 700 °C for 1.5 h. At the moment, Mg was added into the crucible with aluminum foil coated and stirred well. Then the Graphene mixed with aluminum powder was added into

Effects of Graphene Content and Aging Process … Table 1 Compositions of the studied A356.2 alloy (wt%)

Table 2 Aging process of A356.2 alloy with graphene addition

309

Ingredient

Si

Mg

Fe

Other element

Al

wt%

7.00

0.40

1.18 × 106 ; c σ a  560 MPa, N f > 6.22 × 105

several aspects:first of all, the cracks initiate from the pores. Second, the cracks pass through the coating layer. The last, the cracks grow into the CMSX-2 substrate. For the MCrAlY-coated IN738LC and CM247LC specimens, the origin and growth pattern of the cracks are the same [11, 12]. For TBCs coated single crystal superalloy DD6, the cross-sectional structures after fatigue tests are shown in Fig. 4. It is shown that with the increase of stress, holes and cracks decrease because of shorter fracture time. The fracture surfaces of TBCs coated alloy are shown in Fig. 5. The evolution process can be visualized in Figs. 4 and 5. A crack seems to initiate from pores in the bond coat, and then propagates perpendicularly through the bond coat, subsequently grow into the substrate. Such behaviors have been documented for the CoCrAlY-coated superalloys CMSX-2 [11]. There are two main types of fracture surfaces under different stresses: one consists of several planes and the other consists of one plane. DD6 single crystal superalloy owns face-centered cubic structure, and {111} is its easy slipping systems, the crack will grow on {111} octahedral planes. Figure 5a and c show the quasi-cleavage fracture along {111} planes [14, 20]. Figure 6 shows the morphology of the characteristic fatigue striations which are the evidences of fatigue rupture and usually emerged at crack propagation stage. Figure 7 shows the morphology of stepwise characteristic they usually emerged at the final stage of crack propagation, the small planes are {111} octahedral planes. TEM was used to observe the fatigue specimens, due to the nonuniform plastic deformation, the dislocations distribution is not uniform. Dislocation distribution

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Fig. 5 Fracture surface of TBC coated specimens after fatigue tests at 900 °C: a, b σ a  440 MPa, N f  3.35 × 106 ; c, d σ a  500 MPa, N f  1.13 × 106

(b)

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Fig. 6 Fatigue striations on fracture surfaces of specimens: a σ a  500 MPa, N f  1.59 × 106 ; b σ a  520 MPa, N f  9.90 × 105

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Fig. 7 Cleavage steps on fracture surface: a σ a  500 MPa, N f  9.77 × 105 ; b σ a  520 MPa, N f  9.90 × 105

(a)

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Fig. 8 Dislocation configuration of specimens after fatigue tests: a σ a  380 MPa, N f > 1.13 × 107; b σ a  440 MPa, N f  3.35 × 106; c σ a  520 MPa, N f  9.90 × 105

varies with region, and the density of dislocation is large in the area where the deformation is serious. Figure 8a, b, and c show the typical dislocation distribution of the DD6 substrate, the phenomenon of particles sheared by single dislocation and dislocation pairs was not found.

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4 Conclusions 1. The fatigue limits of bared and TBCs coated alloys are 405 and 417 MPa, respectively, the fatigue limit of TBCs coated alloy is increased by 2.96%. 2. There are two main reasons for the difference of fatigue limits between bared and coated alloy at 900 °C: First, the coating can bear part of the stress and make the specimens stronger, second, the residual compressive stresses exists between the coating and substrate. 3. For TBCs coated specimens, the process for eventual failure can be described as a crack which develops from pores in the bond coat and propagates perpendicularly through the bond coat, and finally into the substrate.

References 1. F.A. Zhao, H.Y. Xiao, Z.J. Liu et al., A DFT study of mechanical properties, thermal conductivity and electronic structures of Th-doped Gd2 Zr2 O7 . Acta Mater. 121, 299–309 (2016) 2. A.K. Ray, R.W. Steinbrech, Crack propagation studies of thermal barrier coatings under bending. J. Eur. Ceram. Soc. 19(12), 2097–2109 (1999) 3. J.M. Luo, C.Y. Dai, Y.G. Shen et al., Elasto-plastic characteristics and mechanical properties of as-sprayed 8 mol% yttria-stabilized zirconia coating under nano-scales measured by nanoindentation. Appl. Surf. Sci. 309, 271–277 (2014) 4. J. Wu, H.B. Guo, Y.Z. Gao et al., Microstructure and thermo-physical properties of yttria stabilized zirconia coatings with CMAS deposits. J. Eur. Ceram. Soc. 31(10), 1881–1888 (2011) 5. H.B. Guo, H. Murakami, S. Kuroda, Effect of hollow spherical powder size distribution on porosity and segmentation cracks in thermal barrier coatings. J. Am. Ceram. Soc. 89(12), 3797–3804 (2006) 6. Z.H. Xu, S.M. He, L.M. He et al., Novel thermal barrier coatings based on La2 (Zr0.7Ce0.3)2O7/8YSZ double-ceramic-layer systems deposited by electron beam physical vapor deposition. J. Alloys Compd. 509(11), 4273–4283 (2011) 7. X.L. Chen, Y. Zhao, W.Z. Huang et al., Thermal aging behavior of plasma sprayed LaMgAl11O19 thermal barrier coating. J. Eur. Ceram. Soc. 31(13), 2285–2294 (2011) 8. X. Zhou, Z.H. Xu, R.D. Mu et al., Thermal barrier coatings with a double-layer bond coat on Ni3 Al based single-crystal superalloy. J. Alloys Compd. 591, 41–51 (2014) 9. A.K. Ray, E.S. Dwarakadasa, D.K. Das et al., Fatigue behavior of a thermal barrier coated superalloy at 800°C. Mater. Sci. Eng. A 448(1–2), 294–298 (2007) 10. A.K. Ray, D.K. Das, B. Venkataraman, Characterization of rupture and fatigue resistance of TBC superalloy for combustion liners. Mater. Sci. Eng. A 405(1–2), 194–200 (2005) 11. Y. Itoh, M. Saitoh, Y. Ishiwata, Influence of high-temperature protective coatings on the mechanical properties of nickel-based superalloys. J. Mater. Sci. 34(16), 3957–3966 (1999) 12. Y. Itoh, M. Saitoh, K. Takaki et al., Effect of high-temperature protective coatings on fatigue lives of nickel-based superalloys. Fatigue Fract. Eng. Mater. Struct. 24(12), 843–854 (2001) 13. X.H. Liang, C.G. Deng, M. Liu et al., High cycle fatigue property of NiCoCrAIYTa coating prepared by low pressure plasma spraying on Ni-base single crystal super-alloy at high temperature. J. Therm. Spray Technol. 1(1), 34–38 (2009) 14. Z.X. Shi, J.R. Li, S.Z. Liu et al., High cycle fatigue behavior of the second generation single crystal superalloy DD6. Trans. Nonferrous Met. Soc. China 21(5), 998–1003 (2011) 15. Z.X. Shi, S.Z. Liu, J.R. Li, Rejuvenation heat treatment of the second-generation single-crystal superalloy DD6. Acta Metall. Sin. (Engl. Lett.) 28(10), 1278–1285 (2015)

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16. Y. Liu, J.J. Yu, Y. Xu et al., High cycle fatigue behavior of a single crystal superalloy at elevated temperatures. Mater. Sci. Eng. A. 454–455, 357–366 (2007) 17. J.M. Dong, J.R. Li, R.D. Mu et al., Effect of high temperature heat treatment on elements interdiffusion behavior and stress rupture characteristics of DD6 single crystal superalloy with thermal barrier coatings. J. Mater. Eng. 6, 51–55 (2014) 18. A.K. Ray, B. Goswami, M.P. Singh et al., Characterization of bond coat in a thermal barrier coated superalloy used in combustor liners of aero engines. Mater. Charact. 57(3), 199–209 (2006) 19. A.K. Ray, J.D. Whittenberger, Stress rupture behavior of a thermal barrier coated AE 437A Ni-based superalloy used for aero turbine blades. Mater. Sci. Eng. A 509(1–2), 111–114 (2009) 20. J.Z. Yi, C.J. Torbet, Q. Feng et al., Ultrasonic fatigue of a single crystal Ni-base superalloy at 1000°C. Mater. Sci. Eng. A 443(1–2), 142–149 (2007)

Microstructure and Room-Temperature Fracture Toughness of Nb–Ti–Si In Situ Composite Prepared by Selective Laser Melting Yongwang Kang, Fengwei Guo and Ming Li

Abstract The chemical composition, phase category, and distribution of ternary Nb–Ti–Si in situ composites prepared by arc-melting and selective laser melting (SLM) were investigated by X-ray diffraction and scanning electron microscope equipped with energy dispersive spectroscopy. The room-temperature fracture behaviors were examined by three-point bending tests. SLM would refine the microstructure of Nb–Ti–Si composite, and the finer microstructure was benefit to increase the room-temperature fracture toughness of the corresponding material.

1 Introduction Nbss /Nb-silicide in situ composite materials as potential replacements of Ni-base superalloy in high-temperature application in jet turbine engines were one of the recent attractive research zones, because of their higher melting point, excellent hightemperature strength and low density [1]. Lots of research including the influence of the fabrication methods [2–8] and the role of the chemical composition [9–13] have been performed to obtain the balance of mechanical properties and environment resistance properties for Nbss /Nb-silicide materials. However, until now the improvement of room-temperature fracture toughness is still one of the most challenges to the expected application of this material. Grain size, as one of the factors linking to many properties of materials, has been reported to affect the fracture toughness [14–16]. Selective laser melting (SLM) as a rapid forming technology widely used for metallic materials fabrication may refine the microstructure and improve the mechanical properties comparable to those of forged materials [17, 18]. There are few reports of laser melting deposited Nbss /Nb-silicide in situ composite. In this paper, the microstructures and room-temperature fracture toughness of Nb–Ti–Si ternary

Y. Kang (B) · F. Guo · M. Li Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, 100095 Beijing, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_35

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composite prepared by laser melting deposition were investigated comparing to that by arc-melting method.

2 Experimental A plate of Nb–Ti–Si in situ composite (as-deposit alloy) was fabricated by selective laser melting system with a powder feed nozzle. The raw materials are commercial elements Nb, Ti, Si powder with the particle size of 100–200 μm. Prior to the deposition performance, the elements powders were mixed with a ball mixer. Aside from that, a button ingot of Nb–Ti–Si in situ composite (as-cast alloy) was prepared by arc-melting furnace equipped with nonconsumable electrode and water-cooled copper crucible. The raw materials for as-cast alloy were bulk commercial Nb, Ti, Si. In order to ensure the homogenous of chemical composition, the in situ composite material button was re-melted four times. X-ray diffraction (XRD) analysis was performed on every alloy to characterize phases with Cu-Kα radiation. Microstructures were observed using scanning electronic microscope (SEM). The average chemical compositions of the alloys and the phase distributed in the alloys were identified by energy dispersive microscopy (EDS). The single-edge notched bending (SE(B)) beam specimens with the dimension of 30 mm long, 6 mm wide (W ), 3 mm thick (B) and 3 mm initial notch depth were used to measure the quasi-plane-strain fracture toughness, K Q based on the description in the standard (ASTM E1280-08a). All the detail of the room-temperature fracture toughness testing were mentioned in our previous published paper [19].

3 Results and Discussion 3.1 Microstructure and Chemical Composition XRD spectra of as-deposit alloy and as-cast alloy are shown in Fig. 1. XRD analysis determined that the phase composition of as-deposit alloy was Nbss and Nb3 Si, which was identical to that of as-cast alloy. Generally known, based on the phase diagram of Nb–Ti–Si ternary system [20] Nb3 Si phase as a high-temperature-stable phase was unstable below 1700 °C that with the decrease of temperature Nb3 Si will dissociated to Nbss + Nb5 Si3 as a eutectoid reaction. However, in the two kinds of investigated alloys Nb3 Si was observed at room temperature. This phenomenon indicated that the eutectoid reaction of Nb3 Si to Nbss + Nb5 Si3 did not perform in the solidification process. The reason is related to the low disintegration rate of Nb3 Si that if the cooling rate is high enough there will not be enough time to initiate or finish the eutectoid reaction.

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Nbss

Intensity

Nb3Si

as-deposit

as-cast 10

20

30

40

50

60

70

80

90

2theta /deg. Fig. 1 XRD spectra of as-deposit alloy and as-cast alloy Table 1 The actual compositions of the investigated materials

Alloy

Composition (at%) Nb

Ti

Si

As-deposit

67.23

19.57

13.20

As-cast

67.21

20.76

12.03

Table 1 lists the average chemical compositions of the two kinds of alloys measured with EDS. The chemical compositions are similar, while the content of Si, which mainly concentrated in silicide, in as-deposit alloy was a bit higher than that in as-cast alloy. That means the volume fraction of silicide in as-deposit alloy should be higher than that in as-cast alloy. The microstructure of the two kinds of alloy were shown in Fig. 2. From Fig. 2, the morphology of the as-deposit alloy differed to that of as-cast alloy obviously. The microstructure of as-deposit alloy is very uniform that short-bar Nbss with the size of ~20 μm in Nb-rich Nb3 Si substrate (Fig. 2a). Although the element particle of Nb, Ti and Si did not be detected in the as-deposit alloy, which indicated that the alloying reaction between them was sufficient, at the phase boundary of Nbss and Nb-rich Nb3 Si, the black phase was observed which is characterized to Ti-rich Nb3 Si by EDS. All of the above suggested that because the higher cooling rate in laser melting deposition process, Nbss would not grow to big dimension and Ti element would not homogenized its dislocation in Nb3 Si phase. In as-cast alloy, the microstructure was much coarser than that of as-deposit alloy (Fig. 2b), while Nbss was not uniform with the most size of >50 μm. And between the phase boundary, a little bit of Ti-rich Nb3 Si was found in as-cast alloy. The

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(a)

(b) Nb-rich Nb3Si Nbss

Nb-rich Nb3Si Ti-rich Nb3Si

Nbss

Ti-rich Nb3Si

50 μm

Fig. 2 Microstructures (BSE images) of a as-deposit alloy and b as-cast alloy

phenomenon indicated that the diffusion rates of the elements were much higher in casting process. And with the quantitative image analysis for Fig. 2, the volume fraction of Nbss for as-deposit alloy was about 58.4%, and that for as-cast alloy was about 64.5%, which was consistent to the result of chemical composition analysis.

3.2 Room-Temperature Fracture Toughness Figure 3 compares the load–displacement curves of as-deposit and as-cast alloys. It was seen that the curves of both alloys remained nearly linear up to maximum load with a similar slope ratio. The maximum load for as-cast alloy was 245.86 MPa, while it was 304.59 MPa for as-deposit alloy. The room-temperature fracture toughness (K Q ) was calculated with the method described in Ref. [19]. The results of calculations showed that K Q for as-deposit alloy was 13.09 MPa m1/2 , which is about 23% higher than that for as-cast alloy. In addition, the displacement of as-deposit alloy is higher than that of as-cast alloy. In the post-maximum-load period, the load for as-cast alloy was decreasing gradually, whereas the load for as-deposit alloy dropped suddenly. The three-point bending test results illustrated that the ductility of as-deposit alloy was higher than that of as-cast alloy. However, from the microstructure and chemical composition analysis, the volume of the ductile phase, Nbss in as-cast alloy was higher, which should increase the ductile of the alloy. What is the reason that the as-deposit alloy was much ductile? That should come from the uniform and refined microstructure. Generally known, based on Hall–Petch equation the fine-grain materials is stronger and more ductile than one that was coarse grained. Figure 4 gave fracture surface morphology of both alloys after three-point bending tests with two different magnifications. The fracture surface of as-deposit alloy was relatively uniform (Fig. 4a). However, although the fracture surface of as-cast alloy was also even (Fig. 4c), between the wider tearing ridges much larger quasi-flat areas were observed than that in as-deposit alloy. The difference of fracture surface morphology would be from the difference of microstructure of the alloys. The tearing

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300

Load /N

250 200 150 100 50 0 0.00

0.02

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Displacement /mm Fig. 3 Load–displacement curves of as-deposit alloy and as-cast alloy from three-point bending tests at room temperature

ridges were the morphology of deformed Nbss phase (white arrows in Fig. 4b and d) and the quasi-flat areas were the cleavage facets on the surface fracture, which were corresponding to Nb3 Si phases (black arrows in Fig. 4b, d). The observation for fracture surface indicated that the rupture mechanism of both alloys was the brittle cleavage mode. Combined the calculation results of K Q and fracture surface characterization to microstructure observation, it was suggested that finer microstructure of Nb–Ti–Si in situ composite was benefit to improve its room-temperature fracture toughness. And in as-deposit alloy, much more Ti-rich Nb3 Si (Fig. 2) might affect the room-temperature fracture toughness of as-deposit alloy, which will be investigated furtherly. However, Nb3 Si phase is high-temperature stable phase and is able to be removed by heat treatment, which will improve the room-temperature ductility of Nb-Si alloy.

4 Conclusion Nb–Ti–Si ternary in situ composites were prepared by arc-melting and laser melting deposition methods. The microstructures and room-temperature fracture toughness analysis suggested the following conclusions: (1) Laser melting deposition method was successfully used to prepare Nb–Ti–Si in situ composite. The phase composition is similar to the corresponding alloy prepared by arc-melting method, whereas the microstructure was finer than that formed in arc-melting process.

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(a)

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(d)

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Fig. 4 Fracture surface morphology of a, b as-deposit alloy and c, d as-cast alloy with two different magnification

(2) The room-temperature facture toughness of as-deposit Nb–Ti–Si composite was about 23% higher than that of as-cast composite. The improvement would be induced by the finer microstructure, although the rupture mechanism is still brittle cleavage mode. Acknowledgements This work was financially supported by National Key R&D Program of China (No. 2017YFB0702904).

References 1. J.C. Zhao, J.H. Westbrook, MRS Bull., (September) 622–627 (2003) 2. B.P. Bewlay, M.R. Jackson, J.-C. Zhao, P.R. Subramanian, Metall. Mater. Trans. A 34, 2043–2052 (2003) 3. X.J. Li, H.F. Chen, J.B. Sha, H. Zhang, Mater. Sci. Eng. A 527, 6140–6152 (2010) 4. Y.W. Kang, S.Y. Qu, J.X. Song, Y.F. Han, Acta Metall. Sin. 44, 593–597 (2008) 5. Y.X. Tian, J.T. Guo, G.M. Cheng, L.Y. Sheng, L.Z. Zhou, L.L. He, H.Q. Ye, Mater. Des. 30, 2274–2277 (2009) 6. L.N. Jia, X.J. Li, J.B. Sha, H. Zhang, Rare Metal Mater. Eng. 39, 1475–1479 (2010) 7. R. Dicks, F. Wang, X.H. Wu, J. Mater. Process. Technol. 209, 1752–1757 (2009) 8. Y.W. Kang, S.Y. Qu, Y.F. Han, J.X. Song, D.Z. Tang, Mater. Sci. Forum 561–565, 423–426 (2007) 9. L. Jia, X.P. Guo, Rare Metal Mater. Eng. 36, 1304–1308 (2007) 10. J.T. Guo, Y.X. Tian, L.Y. Sheng, L.Z. Zhou, H.Q. Ye, Int. J. Mater. Res. 99, 1275–1279 (2008)

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Simulation and Experimental Study on the Directional Solidification Process of a Single-Crystal Superalloy Plate Casting Runnan Wang, Qingyan Xu and Baicheng Liu

Abstract The directional solidification (DS) process and as-cast microstructure of a single-crystal (SX) plate casting were investigated through simulations and experiments. Simulations were conducted to optimize the directional solidification process. The mushy zone showed flat and concave morphology at the withdrawing rate of 3 and 6 mm/min, respectively. Based on the simulation results, the directional solidification experiment was performed at the withdrawing rate of 3 mm/min to achieve a preferable microstructure. On the basis of the start block, many newly formed grains appeared due to the chilling effect of the chilling copper plate. With the increase of location height, these grains continuously coarsened and competed, and the grain density gradually decreased. Only a few grains entered into the spiral selector, and only one grain with the orientation close to [001] preferential orientation could get through it to form the SX structure of the plate casting. From the lower to the upper part of the SX plate casting, the primary dendritic arm spacing first decreased from 467.7 to 435.22 µm, and then increased to 565.81 µm; the secondary dendritic arm spacing increased by 11.4%, from 181.39 to 202.13 µm. The γ  phases in the dendritic arms showed a cubical morphology, and some micropores appeared near the eutectic particles. The simulated grain structure evolutions conformed well with the experiments.

1 Introduction Single-crystal Ni-base superalloys play crucial roles in the turbine blades of aeroengine and industrial gas turbine (IGT) due to their extraordinary high-temperature mechanical properties and corrosion resistance [1]. High-rate solidification (HRS) [2] and liquid–metal cooling (LMC) [3, 4] are two main methods widely used for SX component manufacturing. The simulation method, including the temperature distribution, flow field, and structure field can significantly enhance the capability R. Wang · Q. Xu (B) · B. Liu Key Laboratory for Advanced Materials Processing Technology (MOE), School of Materials Science and Engineering, Tsinghua University, Beijing 100084, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_36

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of industrial manufacturing process. Miller et al. [5] investigated the influence of non-axial thermal field on the lateral growth of dendrite during the directional solidification (DS) process and established the criterion for it. Zhang et al. [6] proposed a multi-scale simulation scheme for the simulation of grain and dendrites evolution during the DS process and successfully simulated these in DD6 superalloy. Ma et al. [7] proposed a model for the prediction of the primary spacing. In order to refine the as-cast microstructure, Wang et al. combined the thin mould shell technology with a downward directional solidification process and obtained the castings with finer dendritic microstructure compared to the conventional process [8–12]. The defect formation in the SX components was also widely investigated through experiments and simulations. Yang et al. [13] studied the stray grain formation in the start block of spiral selector during SX seed melting back process and found that the withdrawal velocity, thermal gradient, crystallographic orientation, and alloy properties interactively influenced the crystal growth. Moreover, stray grain is a main defect observed frequently in the SX castings, which is induced by unstability of thermal field [14]. This defect usually occurs in the platform, and the isolated undercooling regions play the crucial roles [15, 16], as well as the geometrical structure of the component [17]. Moreover, the defects such as freckles [18–20], sliver defects [21], surfaced eutectics [22], and recrystallization [23–25] were also widely investigated through experiments and simulations. In this work, the simulations of thermal fields during the DS process with the withdrawing rate of 3 and 6 mm/min were conducted first to obtain a preferable process, and then the DS experiment was carried out. A cellular automaton (CA) method based on KGT crystal growth model was proposed for the simulation of grain evolution during DS process. The grain evolution in the start block of the spiral selector was compared with the simulated structural field, and the simulation results conform well to the experiments. The variation of primary and secondary dendritic arm spacing in different transverse sections of the plate casting was calculated, and the microstructure was characterized.

2 Mathematical Models 2.1 Temperature Field Model The heat transfer process is complicated during directional solidification, including: (1) the radiation between mould shell, casting, and furnace, (2) the heat conduction between chill, casting, core and shell, and (3) the internal heat conduction within the casting, core, and shell. Finite difference method was used for the calculation. The energy conservation model of temperature field follows:        ∂T ∂ ∂T ∂ ∂T ∂T ∂ λ + λ + λ + Q net (1) ρc  ∂t ∂x ∂x ∂y ∂y ∂z ∂z

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where T is the temperature, ρ the density, c the specific heat, λ the heat conductivity, and Qnet the heat flux in the interface elements. A ray-tracking method is employed for the calculation of the radiation heat transfer. The model based on the Stefan–Boltzmann law can be described as: Q net  σ

  ϕn T 4 − Tn4

N  n1

1−ε ε

+

(1−εn )A εn An

+1

(2)

where σ is the Stefan–Boltzmann constant, N the ray line number, ϕ n the energy coefficient of ray line, T the temperature at the start of the ray line, T n the temperature at the end of the ray line, ε the radiation coefficient, A, and An the surface area of the radiation elements.

2.2 Microstructure Model A CA model coupled the temperature field was established for the simulation of grain nucleation and growth during DS process. The nucleation rate is related to the undercooling, which follows:   Ns ∂N ( T − TN )2 √ exp − ∂( T ) 2( Tσ )2 2π Tσ

(3)

where N is the nucleus density, T the average undercooling, N s the maximum nucleus density, T σ the standard deviation of the Gauss distribution, and TN the average nucleation undercooling. The grain density can be described as: T N ( T )  0

  dN d T   d( T )

(4)

The velocity of crystal growth follows KGT model, which can be expressed as: vt  α T 2 + β T 3

(5)

where α and β are the constants which are related to the materials.

3 Experimental Procedures The SX plate castings, 190 mm in length, 50 mm in width and 10 mm in thickness were directionally solidified in an industrial Bridgman furnace with DD407 Ni-base

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Table 1 Nominal chemical composition of the DD407 superalloy Element

Cr

Co

Mo

W

Al

Ti

Ta

Ni

wt%

8

5.5

2.25

5

6

2

3.5

Bal

superalloy, whose nominal chemical composition is presented in Table 1. The liquidus and solidus temperature is 1366 and 1336 °C, respectively. The spiral selector, consisting of start block and spiral part, was added to the very base of the wax pattern to form the SX structure of casting. The whole plate was immerged in the solution of 90% HCl and 10% C2 H5 OH to show the macroscopic grain structure. The microstructural evolutions in the transverse sections of starter block (A1–A3 in Fig. 1) were characterized by Zeiss AM10 optical microscope (OM). Small specimens numbered S1–S4 were cut from the plate casting in upper, middle, lower, and side part of the plate (Fig. 1) using electrical discharge machining (EDM). Their transverse and longitudinal sections were mechanically polished and then etched in the Marble’s reagent (20 g CuSO4 , 100 ml HCl and 100 ml H2 O) to observe the dendritic morphology. The primary (PDAS)

and secondary (SDAS) dendritic arm spacing were calculated by: λPDAS 

A n1

and λSDAS  nL2 , respectively, where A denotes the counting area, n1 the number of dendritic cores in the transverse section, L the length of primary dendritic arm and n2 the number of secondary dendritic arms in the longitudinal section. The samples for scanning electron microscope (SEM) characterizations were electrochemically etched using a solution consisting of 150 ml H3 PO4 , 10 ml H2 SO4 , 15 g CrO3 , and 10 ml H2 O. A Zeiss Merlin VP compact field emission gun (FEG)-SEM was employed.

4 Results and Discussion 4.1 Temperature Distribution The temperature distributions during DS process are closely related to the defect formation and microstructure refinement of SX castings. The thermal field of the SX plate casting during the DS process with the withdrawing rate of 6 mm/min is shown in Fig. 2. The mushy zone shows a concave morphology at the lower part due to the rapid withdrawing rate. The heat dissipation on the side part of the plate casting is better than the center part, which can lead to a lateral dendritic grow. Then the mushy zone gradually evolves to an approximately flat morphology. If the withdrawing rate decreases to 3 mm/min, the mushy zone shows a nearly flat morphology, which is better for crystal growth and defect inhibition (Fig. 3). Hence, this withdrawing rate is chosen for the DS experiment.

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Fig. 1 Structural evolution of the SX plate casting and the specified locations for microstructure characterization in specimens

4.2 Spiral Selector The macroscale structural evolution is shown in Fig. 1, as well as the magnified photograph of the spiral selector. The crystal evolution can be clearly observed. The grain morphologies in the transverse sections (A1–A3) of the start block are shown in Fig. 4. Plenty of grains rapidly nucleate in the surface between the casting and cooling chill due to the chilling effect (Fig. 4a). With the increase of the location height, the newly formed grains gradually coarsen and compete along the heat flow direction, and the grain density gradually decreases. Only a few grains enter into the spiral selector, and only one grain with the orientation close to [001] preferential orientation can get through it to form the SX structure of the plate casting. The CA simulation results of the grain structural evolution are shown in Fig. 5. The grain morphology and the variation of grain density conform well with the experimental results (Figs. 1 and 5).

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Fig. 2 The simulated temperature distribution (a–d) and the mushy zone morphology (f–i) in the SX plate during the directional solidification process with the withdrawing rate of 6 mm/min

Fig. 3 The simulated temperature distribution (a–d) and the mushy zone morphology (f–i) in the SX plate during the directional solidification process with the withdrawing rate of 3 mm/min

4.3 Microstructural Evolution The dendritic evolutions in the transverse and longitudinal sections of specimens S1–S4 are shown in Figs. 6 and 7, respectively. SX dendritic microstructure can be observed, and the direction of dendrite was arranged orderly in the observation surface. The tertiary dendritic arms in Fig. 6d (S4) are more well-developed than the others. The primary and secondary dendritic arm spacing of S1–S4 are shown in Table 2. From the lower to the upper part of the SX plat, the primary spacing

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Fig. 4 Structural evolutions of the chilling grains in different transverse sections of the start block. a section A1, b section A2, and c section A3

Fig. 5 The simulated grain structural evolution from the base of spiral selector to the SX plate

first decreases from 467.7 to 435.22 µm, and then increases to 565.81 µm; the secondary spacing increases by 11.4%, from 181.39 to 202.13 µm. Dendritic growth is a complicated process controlled by cooling rate, temperature gradient, and inclination angle. Under this withdrawing rate, the mushy zone of S1–S3 is very close to a flat morphology. Hence, their PDAS values are close to each other. The fluctuation between them may be induced by the dendritic competition or random factors during dendritic growth process, while the greater PDAS of S4 is introduced by its different thermal condition. In the upper part of the casting, the heat dissipation process is slower due to the heat preservation capability of the pouring system; the mushy zone becomes wider in this part (Fig. 3i), leading to a much greater PDAS in S4.

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Fig. 6 Dendritic evolution in the cross sections of a S1, b S2, c S3, and d S4

Fig. 7 Dendritic evolution in the longitudinal sections of a S1, b S2, c S3, and d S4

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Table 2 PDAS and SDAS of S1–S4 S1

S2

S3

S4

PDAS (µm)

467.70

435.22

463.19

565.81

SDAS (µm)

181.39

196.68

200.68

202.13

As the solidification sequence of the liquid superalloy is different, leading to three different regions in the matrix of SX superalloy: dendritic arm (DA), interdendritic region, and eutectic (EU) as shown in Fig. 8. There are some micropores (MP) near the eutectics after solidification in the as-cast samples. γ (FCC) and γ  (L12) are the phases arranged coherently in the matrix, and the latter one is the main strengthening phases in Ni-base superalloy. In the as-cast samples, the γ  phases exhibit approximately cubical morphology (Fig. 8b), revealing a good microstructure quality of this manufacturing process. The microstructure morphology and composition in different location height are almost the same.

5 Conclusions The directional solidification process and as-cast microstructure were investigated through experiments and simulations in a SX plate casting of DD407 Ni-base superalloy. The following conclusions can be drawn: (1) A model coupling finite difference method and CA method was proposed to simulate the thermal distribution and grain evolution during DS process. The mushy zone transforms from concave to flat morphology if the withdrawing rate decreases from 6 to 3 mm/min.

Fig. 8 Microstructure in the cross section of S1 a DA, EU, and MP near EU, b cubical γ  precipitated phases

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(2) A large number of grains nucleate in the lower part of start block due to the chilling effect; after coarsening and competing, only a few can enter into the spiral selector, and only one grain with the orientation close to [001] can get through it. This grain evolution behavior was successfully simulated. (3) The tertiary dendritic arms in the upper part of the plat casting are more welldeveloped than the lower part due to different heat dissipation conditions. From the lower to the upper part of the SX casting, the primary dendritic arm spacing firstly decreases from 467.7 to 435.22 µm, and then increases to 565.81 µm; the secondary dendritic arm spacing increases by 11.4%, from 181.39 to 202.13 µm. The γ  phases in the dendritic arms show a cubical morphology, and some micropores appear near the eutectic particles. Acknowledgements This research was funded by the National Science and Technology Major Project (2017ZX04014001) and National Key R&D Program of China (2017YFB0701503).

References 1. J.H. Perepezko, The hotter the engine, the better. Science 326(5956), 1068–1069 (2009) 2. R.C. Reed, The superalloys fundamentals and applications (Cambridge University Press, New York, 2006) 3. M.M. Franke, R.M. Hilbinger, A. Lohmüller, R.F. Singer, The effect of liquid metal cooling on thermal gradients in directional solidification of superalloys: thermal analysis. J. Mater. Process. Technol. 213(12). 2081–2088 (2013) 4. J.D. Miller, T.M. Pollock, Stability of dendrite growth during directional solidification in the presence of a non-axial thermal field. Acta Mater. 78, 23–36 (2014) 5. J.D. Miller, L. Yuan, P.D. Lee, T.M. Pollock, Simulation of diffusion-limited lateral growth of dendrites during solidification via liquid metal cooling. Acta Mater. 69, 47–59 (2014) 6. H. Zhang, Q. Xu, B. Liu, Numerical simulation and optimization of directional solidification process of single crystal superalloy casting. Materials 7(3), 1625–1639 (2014) 7. D. Ma, P.R. Sahm, Primary spacing in directional solidification. Metall. Mater. Trans. 29A, 1113–1119 (1998) 8. F. Wang, D.X. Ma, J. Zhang, S. Bogner, A. Bührig-Polaczek, A high thermal gradient directional solidification method for growing superalloy single crystals. J. Mater. Process. Technol. 214(12), 3112–3121 (2014) 9. F. Wang, D. Ma, J. Zhang, S. Bogner, A. Bührig-Polaczek, Solidification behavior of a Ni-based single crystal CMSX-4 superalloy solidified by downward directional solidification process. Mater. Charact. 101, 20–25 (2015) 10. F. Wang, D. Ma, J. Zhang, L. Liu, J. Hong, S. Bogner, A. Bührig-Polaczek, Effect of solidification parameters on the microstructures of superalloy CMSX-6 formed during the downward directional solidification process. J. Cryst. Growth 389, 47–54 (2014) 11. F. Wang, D. Ma, J. Zhang, A. Bührig-Polaczek, Investigation of segregation and density profiles in the mushy zone of CMSX-4 superalloys solidified during downward and upward directional solidification processes. J. Alloy. Compd. 620, 24–30 (2015) 12. F. Wang, D. Ma, J. Zhang, L. Liu, S. Bogner, A. Bührig-Polaczek, Effect of local cooling rates on the microstructures of single crystal CMSX-6 superalloy: a comparative assessment of the bridgman and the downward directional solidification processes. J. Alloy. Compd. 616, 102–109 (2014)

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13. X. Yang, D. Ness, P.D. Lee, N. D’Souza, Simulation of stray grain formation during single crystal seed melt-back and initial withdrawal in the Ni-base superalloy CMSX-4. Mater. Sci. Eng. A413–414, 571–577 (2005) 14. D. Szeliga, K. Kubiak, J. Sieniawski, Control of liquidus isotherm shape during solidification of Ni-based superalloy of single crystal platforms. J. Mater. Process. Technol. 234, 18–26 (2016) 15. X.B. Meng, J.G. Li, Z.Q. Chen, Y.H. Wang, S.Z. Zhu, X.F. Bai, F. Wang, J. Zhang, T. Jin, X.F. Sun, Z.Q. Hu, Effect of platform dimension on the dendrite growth and stray grain formation in a Ni-base single-crystal superalloy. Metall. Mater. Trans. 44A, 1955–1965 (2012) 16. X.B. Meng, J. Li, S. Zhu, H. Du, Z. Yuan, J. Wang, T. Jin, X. Sun, Z. Hu, Method of stray grain inhibition in the platforms with different dimensions during directional solidification of a Ni-base superalloy. Metall. Mater. Trans. 45A, 1230–1237 (2013) 17. R. Wang, X. Yan, Z. Li, Q. Xu, B. Liu, Effect of construction manner of mould cluster on stray grain formation in dummy blade of DD6 superalloy. High Temp. Mater. Processes 36(4), 399–409 (2017) 18. J. Hong, D. Ma, J. Wang, F. Wang, A. Dong, B. Sun, A. Bührig-Polaczek, Geometrical effect of freckle formation on directionally solidified superalloy CM247 LC components. J. Alloys Compd. 648, 1076–1082 (2015) 19. D. Ma, A.B. Polaczed, The influence of surface roughness on freckle formation in directionally solidified superalloy samples. Metall. Mater. Trans. 43B, 671–677 (2012) 20. D. Ma, Q. Wu, B.P. Andreas, Some new observations on freckle formation in directionally solidified superalloy components. Metall. Mater. Trans. 43B, 344–357 (2012) 21. J.W. Aveson, P.A. Tennant, B.J. Foss, B.A. Shollock, H.J. Stone, N. D Souza, On the origin of sliver defects in single crystal investment castings. Acta Mater. 61(14), 5162–5171 (2013) 22. L. Cao, L. Yao, Y. Zhou, T. Jin, X. Sun, Formation of the surface eutectic of a Ni-based single crystal superalloy. J. Mater. Sci. Technol. 33(4), 347–351 (2017) 23. H.N. Mathur, C. Panwisawas, C.N. Jones, R.C. Reed, C.M.F. Rae, Nucleation of recrystallisation in castings of single crystal Ni-based superalloys. Acta Mater. 129, 112–123 (2017) 24. C. Panwisawas, H. Mathur, J. Gebelin, D. Putman, C.M.F. Rae, R.C. Reed, Prediction of recrystallization in investment cast single-crystal superalloys. Acta Mater. 61(1), 51–66 (2013) 25. L.H. Rettberg, T.M. Pollock, Localized recrystallization during creep in nickel-based superalloys GTD444 and René N5. Acta Mater. 73, 287–297 (2014)

Influence of Mo Content on the Microstructure Stability and Stress Rupture Properties of a Single Crystal Superalloy Z. X. Shi, S. Z. Liu, X. G. Wang and J. R. Li

Abstract The single crystal superalloy with 1 and 3% Mo was prepared using the high vacuum induction melting furnace. The effect of Mo content on the structural stability and endurance properties of the alloy was studied. The results showed that the γ  precipitate size reduced and turned uniform and the cube form became regular with increasing Mo content. The γ  directional rafting and topologically close-packed phase precipitation appeared in the 1 and 3% Mo alloy after aged for 1000 h at 1100 °C. The volume fraction of topologically close-packed phase significantly enhanced with rise of Mo percentage. The stress rupture properties at 1100 °C/140 MPa of the alloy depredated with rise of Mo percentage.

1 Introduction Single crystal Ni-base superalloys are typically used for aircraft engine turbine blades on account of their excellent mechanical properties at high temperature [1–3]. Mo has a big role as one of refractory alloying element in improving the performance of the single crystal superalloy [4, 5]. It can reduce the diffusion rate of alloying elements to delay γ  directional rafting [6]. Mo can facilitate to form dense γ /γ  interfacial dislocation network to reduce creep rate [7]. So the single crystal superalloys almost contain Mo element. However, the high-level Mo content made the alloy susceptible to precipitate harmful TCP phases [8–10]. Therefore, it is important to study the role of Mo in a new generation alloy. The present study examined the influence of Mo element on the microstructure and endurance properties of a new single crystal Ni-base superalloy to provide a basis for the alloy design and application.

Z. X. Shi (B) · S. Z. Liu · X. G. Wang · J. R. Li Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing 100095, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_37

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2 Experiment A Ni–Re–Ru–Mo–Cr–Ta–Co–Al–Nb–Hf–Y system single crystal superalloy with different Mo content was prepared using the high vacuum induction melting furnace. The Mo content of two alloys is 1 and 3%, respectively, and other chemical compositions are equivalent. The deviation angle of crystal orientation of each single crystal rod was tested using XRD (X-ray diffractometer). The rods with orientation within 8° deviating from [001] direction were used for the following experiment. The single crystal rod was carried out a standard heat treatment 1335 °C/8 h/air cooling + 1140 °C/4 h/air cooling + 870 °C/24 h/air cooling. The long-term aging was performed for 1000 h at 1100 °C after standard heat treatment. The sample for endurance property test was machined after standard heat treatment. The endurance property was performed at 1100 °C/140 MPa in air. Microstructure observation at different conditions was performed on SEM. The JMatPro software was used to calculate the equilibrium phases of the alloy.

3 Results and Analysis 3.1 Microstructure SEM photograph of heat-treated structure of the superalloy with different Mo content is presented in Fig. 1. It is shown that there was no incipient melting point in the microstructures of as though two alloys were carried out with same heat treatment regime. The coarse γ  phase and (γ + γ  ) eutectic was almost dissolved and the dendrite segregation was greatly alleviated after high-temperature solid solution treatment. They all contained about 65% γ  precipitate with regular cube form and γ phase. The cube edge of γ  particle of the 1 and 3% Mo alloy was about 0.46 and 0.41 μm, respectively. It has the same result with study carried out by Wang [11]. The γ  precipitate size reduced and turned uniform and the cube form became regular with increasing Mo content. Therefore, the γ phase passages of 3% Mo superalloy are much straighter than those of 1% Mo superalloy. SEM photograph of long-term heat-treated structure of the superalloy with different Mo content is illustrated in Fig. 2. γ  directional rafting appeared in the longterm heat-treated process. The directed diffusion of alloy elements at high temperatures resulted in the directional rafting of the γ  phase. There is a lattice mismatch between γ matrix and γ  precipitate as they have a little different lattice parameter. The common-lattice stress as a result of the lattice mismatch can induce γ  phase directional rafting [12]. The acicular topologically close-packed phases precipitated in two alloys after aged 1000 h at 1100 °C, but the amount of TCP phase increased with increase Mo content. It can be seen from Fig. 2 that topologically close-packed phases nucleated and developed along certain plane and orientation in two alloys. The topologically close-

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Fig. 1 SEM photograph of heat-treated microstructure of the superalloys with different Mo content: a 1.0% Mo; b 3% Mo

Fig. 2 SEM photograph of long-term heat-treated structure of the superalloys with different Mo content: a 1.0%; b 3%

packed phase has a particularly complicated crystalline structure and the single cell dimension is much bigger than that of γ matrix or γ  precipitate. So it is very difficult to precipitate topologically close-packed phase in these alloys [13]. The topologically close-packed phase preferentially forms on the dense plane and grows along the dense orientation to minimize the energy of the system if they precipitate [14]. Therefore, it nucleated and grew along certain plane and orientation and appear obvious direction relations with the matrix. The structure stabilization is one of the important indexes for the alloy served at high temperatures [15]. A larger volume fraction of topologically close-packed phase of the single crystal superalloy precipitated as rise of Mo percentage. It can be concluded that much more Mo percentage added in the alloy containing high content of high melting point alloying elements can reduced structure stabilization at elevated temperature. The precipitation of topologically close-packed phase in superalloy is usually ascribed to the refractory elements oversaturation in the γ matrix [16]. As a γ matrix form element, the partition ratio of Mo can increase with increase its content. Moreover, increase of Mo content raised the distribution ratio of Cr, Re, and W [9, 10]. So increase of Mo content rises the percentage of Cr, Mo, W, and Re in the

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Fig. 3 The equilibrium phases forming sequence of the alloy: a 1% Mo b 3% Mo; c effect of Mo on the phase volume at 1100 °C

γ matrix. Therefore, the oversaturation degrees of topologically close-packed phase form element in the γ matrix are raised, which makes 3% Mo alloy precipitate much more topologically close-packed phase than that the 1% Mo alloy. It has the same result with study investigated by Zhang [6]. The influence of Mo percentage on equilibrium phases forming sequence is studied using the JMatPro calculation procedure and corresponding database. The phases forming sequence of the alloy containing different Mo percentage as the temperature going down is illustrated in Fig. 3. The same phases, carbide, γ  , γ and topologically close-packed phase precipitate in two alloy. However, the 3% Mo alloy precipitates much more topologically close-packed phase than the 1% Mo. It has the same result with the above experimental observations.

3.2 Stress Rupture Properties The endurance properties of the two alloys at 1100 °C/140 MPa illustrated in Table 1. Each number is the mean value of 3 times experiments. It indicates that with increase Mo content, the stress rupture life, elongation, and reduction of area decreased, showing similar result reported in other alloy [9].

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Table 1 The endurance properties of the two alloys at 1100 °C/140 MPa Specimens

t (h)

δ (%)

ψ (%)

1% Mo

382.8

29.8

40.5

3% Mo

317.3

18.2

29.3

(a)

(b)

σ

σ Fig. 4 SEM photograph near the fracture surface of ruptured samples containing different Mo content: a 1% Mo; b 3% Mo

As Mo is a strong solid solution strengthener in the superalloy, it is able to enhance the endurance properties of the alloy [4, 5]. The degradation of endurance properties is possibility ascribed to the microstructure deviation increase of Mo percentage. Figure 4 shows the SEM photograph near the fracture of the ruptured samples containing different Mo content. It is shown that the raft structure has been formed for all the specimens. The samples exhibit the precipitation of topologically closepacked phase in 1 and 3% Mo alloy. Moreover, the 3% Mo alloy precipitate much more topologically close-packed phase than that the 1% Mo alloy, which is in good agreement with experiment carried out in long-term aging. The microcrack can generate at the interfaces between topologically close-packed phase and matrix because concentration of stress [10]. Therefore, the crack had formed on the topologically close-packed phase in the 3% Mo alloy. It indicates that the decrease of endurance properties with increase of Mo percentage can be attributed to the increase of the volume fraction of topologically close-packed phase.

4 Conclusions (1) The γ  precipitate size reduced and the cube form became regular with increasing Mo content. (2) Directional rafting of the γ  phase and topologically close-packed phase precipitation appeared in the 1 and 3% Mo alloys after aged for 1000 h at 1100 °C.

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The volume fraction of topologically close-packed phase significantly increased with rise of Mo percentage. (3) The stress rupture properties at 1100 °C/140 MPa of the alloys decreased with rise of Mo percentage.

References 1. X. Wu, A. Dlouhy, Y.M. Eggeler, E. Spiecker, A. Kostka, C. Somsen, G. Eggeler, On the nucleation of planar faults during low temperature and high stress creep of single crystal Nibase superalloys. Acta Mater. 144, 642–655 (2018) 2. K. Arora, K. Kishida, K. Tanaka, H. Inui, Effects of lattice misfit on plastic deformation behavior of single-crystalline micropillars of Ni-based superalloys. Acta Mater. 138, 119–130 (2017) 3. B. Dubiel, I. Kalemba-Rec, A. Kruk, T. Moskalewicz, P. Indyka, S. Kac, A. Radziszewska, A. Kopia, K. Berent, M. Gajewska, Influence of high-temperature annealing on morphological and compositional changes of phases in Ni-base single crystal superalloy. Mater. Charact. 131, 266–276 (2017) 4. L. Qin, Y.L. Pei, S.S. Li, X.B. Zhao, S.K. Gong, H.B. Xu, Role of volatilization of molybdenum oxides during the cyclic oxidation of high-Mo containing Ni-based single crystal superalloys. Corros. Sci. 129, 192–204 (2017) 5. Y. Ru, S.S. Li, Y.L. Pei, J. Zhou, S.K. Gong, H.B. Xu, Interdendritic Mo homogenization and sub-solidus melting during solution treatment in the Mo-strengthening single crystal superalloys. J. Alloy. Compd. 662, 431–435 (2016) 6. J. Zhang, J.G. Li, T. Jin, X.F. Sun, Z.Q. Hu, Effect of Mo concentration on creep properties of a single crystal nickel-base superalloy. Mater. Sci. Eng., A 527, 3051–3056 (2010) 7. J.X. Zhang, T. Murakumo, H. Harada, Y. Koizumi, T. Kobayashi, Creep deformation mechanisms in some modern single-crystal superalloys, in Superalloys 2004, ed. by K.A. Green, T.M. Pollok, H. Harada, T.W. Howson, R.C. Reed, J.J. Schirra, S. Walston (TMS, Pennsylvania, 2004), pp. 189–195 8. X.G. Liu, L. Wang, L.H. Lou, J. Zhang, Effect of Mo addition on microstructural characteristics in a Re-containing single crystal superalloy. J. Mater. Sci. Technol. 31(2), 143–147 (2015) 9. P.P. Hu, J.Y. Chen, Q. Feng, Y.H. Chen, L.M. Cao, X.H. Li, Effects of Mo on the microstructure and stress rupture property of Ni-based single crystal superalloys. Chin. J. Nonferrous Met. 21, 332–340 (2011) 10. W.Y. Ma, Y.F. Han, S.S. Li, Y.R. Zheng, S.K. Gong, Effect of Mo content on the microstructure and stress rupture of a Ni base single crystal superalloy. Acta Metall. Sin. 42(11), 1191–1196 (2006) 11. B. Wang, J. Zhang, T.W. Huang, H.J. Su, Z.R. Li, L. Liu, H.Z. Fu, Influence of W, Re, Cr, and Mo on microstructural stability of the third generation Ni-based single crystal superalloys. J. Mater. Res. 31(21), 3381–3389 (2016) 12. J. Coakley, E.A. Lass, D. Ma, M. Frost, H.J. Stone, D.N. Seidman, D.C. Dunand, Lattice parameter misfit evolution during creep of a cobalt-based superalloy single crystal with cuboidal and rafted gamma-prime microstructures. Acta Mater. 136, 118–125 (2017) 13. Z. Shi, J. Li, S. Liu, Effects of Ru on the microstructure and phase stability of a single crystal superalloy. Int J Miner, Metall Mater 19, 1004–1009 (2012) 14. Z. Shi, S. Liu, J. Li, Effects of Cr content on microstructure and mechanical properties of single crystal superalloy. Trans. Nonferrous Met Soc Chin 25, 776–782 (2015) 15. J.B. Graverend, J. Cormier, P. Caron, S. Kruch, F. Gallerneau, J. Mendez, Numerical simulation of γ/γ microstructural evolutions induced by TCP-phase in the MC2 nickel base single crystal superalloy. Mater. Sci. Eng., A 528, 2620–2634 (2011) 16. R.A. Hobbs, L. Zhang, C.M.F. Rae, S. Tin, TCP suppression in a ruthenium-bearing singlecrystal nickel-based superalloy. JOM 60(7), 37–42 (2008)

Preparation and Compressive Properties of Advanced Pore Morphology (APM) Foam Elements Yanli Wang, Lucai Wang, Hong Xu and Qiaoyu Guo

Abstract The Advanced Pore Morphology (APM) foam elements were prepared by modified Powder-Compacting Foaming (PCF) method. APM aluminum foam elements have a nearly-spherical surface with closed-cell porous structure and integral skin, with an average diameter of about 15 mm. Their average pore size is 1.879 mm and average equivalent circle circularity is 0.8. Two factors (heating speed and foaming time) played important roles in the preparation of APM foamed aluminum. The deformation of APM foam element exhibits plastic feature under quasi-static compression tests, the elastic region of APM foam aluminum elements is short, and the plastic deformation is not homogeneous.

1 Introduction Aluminum foams are a significant embranchment of metal foams, which is prepared by adding foaming additives into aluminum or aluminum alloy. They are composed of aluminum matrix skeleton and pore phase, and usually used as a core of sandwich structure or a filler of hollow structure for several multifunctional construction in engineering, thus exhibit many attractive properties, such as lightweight, high specific strength, heat insulation, damping, and electromagnetic shielding and so on [1–3]. The important limitation of conventional aluminum foams materials is their stochastic geometry pore, caused physical and mechanical properties inconsistencies [4]. Additionally, their high manufacturing cost also becomes a hindrance to the multiscale industrial application. There are some technical problems in the preparation of large composite aluminum foam parts, such as the temperature field is difficult to be controlled accurately, which leads to the decrease of the pore structure uniformity of the foam core, and it is difficult to guarantee the filling effect of the special-shaped composite parts [5, 6]. Metallic hollow sphere (MHS) structure [7, 8] and APM Y. Wang (B) · H. Xu North University of China, Taiyuan 030051, Shanxi Province, China e-mail: [email protected] L. Wang · Q. Guo Taiyuan University of Science and Technology, Taiyuan 030024, Shanxi Province, China © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_38

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3mm

Fig. 1 APM foam elements

aluminum foam elements [9–13] helped to overcome the technological problems related to the control of cellular structure irregularity because these methods have easy replicability, and it is more homogeneous of the cellular structure and pore distribution. Extremely, hollow spheres or granular porous materials are selected as the appropriate basic materials. The functional gradient core materials and structures have been developed to increase the stiffness and energy absorption for their more convenient and flexible use in different engineering applications [14]. APM foam elements have been developed at Fraunhofer IFAM Bremen in Germany, and their fabrication was derived from the CONFORM process [9–11]. APM foam elements (Fig. 1) have a spherical outer skin and interconnected closed-cell porous structure, usually 3–15 mm in diameter with density varying from 500 to 1000 kg/m3 .

2 Manufacturing Process of APM Foam Elements by Semi-constrained Method The foamable precursor is made up of metal powder (Al (89.2wt%) and Si (10wt%)) and a blowing agent ZrH2 (0.8wt%), named as the foamable precursor. The wireshaped precursor material (diameter  Ø10 mm) is then cut into small granulates (length  10 mm), and putted in the half-hollow stainless steel ball, which are then expanded into spherical-like foam elements due to heat reaction of the ZrH2 foaming agent at the three belt furnace temperature of 750, 780, and 800 °C. So, the sizes of APM foam elements have been manufactured with diameters of 15 mm and average density of 840 ± 40 kg/m3 . The temperature of the sample is determined by the preset furnace temperature. The higher the preset furnace temperature is, the faster the heating rate is, and the higher the heating rate is. Table 1 exhibits the average heating rates at different preset furnace temperatures.

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Table 1 Heating rate of the precursor refers to different furnace temperature Furnace temperature (°C)

750

780

800

Heating rate (°C/s)

2.50

2.85

3.02

(a)

(b)

Time (s)

Fig. 2 a Samples temperature and expansion height; b Cross-section micrographs of foam at maximum expansion, at the furnace temperature 750 °C

Fig. 3 a Samples temperature and expansion height; b Cross-section micrographs of foam at maximum expansion, at the furnace temperature 780 °C

The heating process of the samples is slightly different under different preset furnace temperatures as shown in Figs. 2a, 3a and 4a. At first, there was no significant difference in the heating rate of the samples. When the temperature of the samples approached the melting point, the temperature rising rate gradually slowed down, the plateau areas with different lengths appeared. With continuously absorbing heat and melting of samples, the length of the plateau areas on 800 °C was the shortest. Subsequently, the samples would continue to heat up until they approached the preset furnace temperature and then remained in equilibrium. The important parameters analysis is collected from the experimental curve of Figs. 2a, 3a and 4a, as shown in Table 2. When the foaming process was taking on to about 220 s, the sample temperature just reached the temperature plateau region, the matrix strength was still very high, and hydrogen was difficult to release, so the small granulates did not expand significantly. The precursor was heated and softened

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(a)

(b)

Time (s)

Fig. 4 a Samples temperature and expansion height; b Cross-section micrographs of foam at maximum expansion; at the furnace temperature 800 °C Table 2 Heating and expansion parameters of samples at different furnace temperature Furnace temperature (°C)

Start foaming times (s)

Maximum Maximum Porosity expanExpan- (%) sion sion time (s) height (mm)

Foaming temperature (°C)

Average equivalent diameter (Da ) (mm)

Number of pores

750

260

430

12.3

70

678–743 1.663

45

780

250

330

15.1

77.3

670–705 1.879

40

800

215

332

18.1

82.4

675–747 2.524

20

between 270 and 310 s, and the blowing agent decomposed and produced gas H2 . With processing, the mount of H2 was increased and densely distributed inside the sample, and the sample showed an obvious expansion process. When the time reaches over 330 s, as the cells grow and merge, the sample reaches the maximum expansion point. At this moment, the porosity of the sample reaches the maximum value. If foaming still continued, cells were unstable due to the low viscosity of the matrix and the gravitation, resulting in the cracking of the cells and the overall collapse in the macroscopic view. The cross sections were measured by the binarization processing on image analysis software. Equivalent diameter (Di ), Average equivalent diameter (Da ), Circularity (C i ), and Average equivalent circle circularity (C a ) were calculated by Formula (1–4). Circularity distributions evaluated by equivalent diameter as shown in Figs. 2, 3 and 4b. Equivalent diameter (Di ): Circle diameter of the area of the corresponding hole; Di  2( Ai /π )1/2 Average equivalent diameter (Da ): The ratio of the sum diameter of cross section and the number of holes N;

(1) 

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Ci /N

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Ci on the

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Figure 2 described many scattered small pores in the cross section of the sample, the number of pores is between 45 and 50, the circularity is relatively better, and the equivalent circle diameter is concentrated within 3 mm (Da  1.663, C a  0.808) at the furnace temperature 750 °C. The rate of ZrH2 release accelerated at the furnace temperature 780 °C, the initial formation of the small pores was growing up and contacted each other, the cell wall became thinner and merged. A lot of pores disappeared and formed, the number of pores is about 40 and the average equivalent diameter increased, and the equivalent circle diameter is concentrated within 4 mm (Da  1.879, C a  0.801). The melt viscosity is greatly decreased at the furnace temperature 800 °C, the cell merged quickly, thus the number of pores is reduced to about 20, and the average equivalent circle diameter is further increased (Da  2.524, C a  0.726). The majority of pores in APM foam elements were nonspherical, average equivalent circle circularity C a is 0.8 generally, the pores was random, as no common direction of the primary axis of pore ellipsoids could be identified [5]. Because of the low influence of gravitational and drainage effects on the structural homogeneity of APM foam elements, Lehmhus et al. [15] already confirm this. Above all, when the heating rate was 2.85 °C/s (the preset furnace temperature 780 °C), we will obtain high porosity and circularity when foaming completion time was controlled shorter than 330 s.

3 Compressive Behavior of APM The objective of this experiment was to characterize the behavior of single APM foam elements under quasi-static compressive loading condition. It was carried out on a SANS microcomputer-controlled electronic universal testing machine (model No. CMT5105), by displacement loading method at a constant speed of 3 mm/min. Figure 5 showed the sequential deformation of single APM foam elements. At the initial stage of compression, deformation started at the contact points between loading plates and foam element, the foam ball is obviously flattened from both load introduction points toward the center plane. Subsequently, the middle part of the ball

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Fig. 5 Sequential deformation of single APM foam elements (Ø15 mm) under quasi-static compressive loading conditions Fig. 6 Load–strain curves of APM elements

began to expand and thicken until it was pressed into pancakes. The closed outer skin, which constitutes a density gradient from core to surface, play a very important role in it. Stöbener et al. [11] obtained the approximate result of Ø10 mm APM foam element by experimental compression test. Figure 6 described the quasi-static force–strain response of APM foam elements with Ø15 mm diameter. All five specimens have similar deformation global response. It indicated a typical compressive behavior of aluminum foam, composed of typical three distinct regions: (A) the initial linear elastic region and the slow increase in the stress up to the yield point, (B) followed by plateau stress with a gradual trend of stress increase, and (C) the densification region in which the stress abruptly increases. The plateau and densification values are shown in Table 3. Deformation started at the contact points between loading plates and specimen, the elastic region of APM foam aluminum elements is likely shorter, the plastic deformation is not homogeneous, and the deformation is not limited to the surface of the spherical element, but is continued

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Table 3 Mechanical properties of the Ø15 mm APM foam elements Quasi-static loading

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in the internal foam structure. Vesenjak et al. [16, 17] observed that the yielding originates from the contact area between the APM foam and the rigid support and spreads in a shear band with an angle of 45° toward the lower end of the foam element. Sulong et al. [18] generated an accurate finite element model by µCT techniques to study the internal microstructure deformation, indicated that the weak structures like thin walls and thin struts within the 45° shear zone are likely to experience high levels of plastic deformation and plasticization is not limited to the immediate surface area.

4 Conclusions In the preparation of APM foam, the preset furnace temperature (or the heating rate) and foaming time are important factors. The experimental results showed that high porosity and circularity can be obtained at the preset furnace temperature 780 °C (heating rate 2.85 °C/s). A typical compressive behavior of aluminum foam is observed, and the deformation is not limited to the surface of the spherical element, but continued in the internal foam structure. The yield regions transfer from the contact area in a shear band with an angle of 45° to the lower end of the foam element.

References 1. M. Vesenjak, Z. Ren, Geometrical and mechanical analysis of various types of cellular metals. Ciênc. Tecnol. dos Mater. 28, 9–13 (2016) 2. J. Banhart, Manufacture, characterization and application of cellular metals and metal foams. Prog. Mater. Sci. 46, 559–632 (2001) 3. Moon Sik Han, Jae Ung Cho, Impact damage behavior of sandwich composite with aluminum foam core. Trans. Nonferrous Met. Soc. Chin. 24, 42–46 (2014) 4. J. Baumeister, J. Weise, Applications of aluminum hybrid foam sandwiches in battery housings for electric vehicles. Proc. Mater. Sci. 4, 317–321 (2014) 5. M. Ulbin, M. Borovinšek, Y. Higa, K. Shimojima, Internal structure characterization of AlSi7 and AlSi10 advanced pore morphology (APM) foam elements. Mater. Lett. 136, 416–419 (2014) 6. A. Kovaˇciˇc, Z. Ren, On the porosity of advanced pore morphology structures. Compos. Struct. 158, 235–244 (2016) 7. Luo Xin, The mechanical behavior of thin-walled tube filled with hollow metal spheres. Compos. Struct. 133, 124–130 (2015) 8. O. Friedl, C. Motz, H. Peterlik, S. Puchegger, N. Reger, R. Pippan, Experimental investigation of mechanical properties of metallic hollow sphere structures. Metall. Mater. Trans. B 39(1), 135–146 (2008)

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9. K. Stöbener, J. Baumeister, G. Rausch, M. Rausch, Forming metal foams by simpler methods for cheaper solutions. Met. Powder Rep. 60(1), 12–16 (2005) 10. K. Stöbener, J. Baumeister, G. Rausch, M. Busse, Metal foams with advanced pore morphology (APM). High Temp. Mater. Processes (London) 26(4), 231–238 (2007) 11. K. Stöbener, D. Lehmhus, M. Avalle, L. Peroni, M. Busse, Aluminum foam-polymer hybrid structures (APM aluminum foam) in compression testing. Int. J. Solids Struct. 45, 5627–5641 (2008) 12. K. Stöbener, G. Rausch, Aluminium foam-polymer composites: processing and characteristics. J. Mater. Sci. 44, 1506–1511 (2009) 13. I. Duarte, M. Vesenjak, Compressive performance evaluation of APM (Advanced pore morphology) foam filled tubes. Compos. Struct. 134, 409–420 (2015) 14. J. Hohe, V. Hardenacke, V. Fascio, Y. Girard, J. Baumeister, Numerical and experimental design of graded cellular sandwich cores for multifunctional aerospace applications. Mater. Des. 39, 20–32 (2012) 15. D. Lehmhus, J. Baumeister, L. Stutz, E. Schneider, K. Stöbener, M. Avalle, Mechanical characterisation of particulate aluminum foams—strain-rate, density and matrix alloy versus adhesive effects. Adv. Eng. Mater. 12, 596–603 (2010) 16. M. Vesenjak, M. Borovinšek, T. Fiedler, Y. Higa, Z. Ren, Structural characterisation of advanced pore morphology (APM) foam elements. Mater. Lett. 110, 201–203 (2013) 17. M. Vesenjak, F. Gacnik, L. Krstulovicopara, Mechanical properties of advanced pore morphology foam elements. Mech. Adv. Mater. Struct. 22(5), 359–366 (2015) 18. M.A. Sulong, M. Vesenjak, I.V. Belova, G.E. Murch, T. Fiedler, Compressive properties of advanced pore morphology (APM) foam elements. Mater. Sci. Eng., A 607, 498–504 (2014)

Effect of Temperature and Strain Amplitude on Low Cycle Fatigue Behaviour of DD11 Yuanyuan Guo, Yunsong Zhao, Jian Zhang, Yanfei Liu, Zhenyu Yang, Jiang Hua and Yushi Luo

Abstract The fatigue behaviours, including the fatigue cycle life, cyclic stress response, microstructure evolution and the fatigue fracture morphology of DD11 alloy with [0 0 1] orientation under the isothermal low cycle fatigue condition at 760 and 980 °C were investigated. The fatigue life decreased significantly with increasing total strain amplitude, but did not decrease with increasing temperature. The cyclic softening was related to the climb and cross-slip of dislocations. The dislocations which cut through the γ  phase caused cyclic hardening. The fracture mode included shear fracture and normal fracture. This work is beneficial to build the relationship between microstructure and cyclic stress response behaviour.

1 Introduction Nickel-based single crystal superalloys have been widely performed as blade materials in the gas turbines since their excellent mechanical and corrosion resistance at the elevated temperature [1–4]. The blade suffered repeated thermal stresses and the low cycle fatigue was the major failure mode [5–7]. Plenty of research results indicated that the temperature [8, 9], loading rate [10], waveform [11] and strain amplitude [12] were the main factors to influence the fatigue life and fracture mechanisms. The cyclic stress response behaviour was different as the temperature and strain amplitude changing. At the low and intermediate temperature, the deformation mechanism was planer slip. The fatigue crack propagates at the crystallographic plane which was oblique to the loading direction. As the temperature increases, the fracture surface of fatigue was more perpendicular to the loading axis, in addition, the wavy slip was the main deformation. Y. Guo (B) · Y. Zhao · J. Zhang · Z. Yang · J. Hua · Y. Luo Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing 100095, China e-mail: [email protected] Y. Liu Sichuan Gas Turbine Research Establishment, Sichuan 610500, China © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_39

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The present study is devoted to the low cycle fatigue (LCF) behaviour of a secondgeneration nickel-based superalloy single crystal DD11, an alloy developed by Science and Technology on Advanced High Temperature Structural Materials, Beijing Institute of Aeronautical Materials. The influence of temperature and strain amplitude was investigated. The LCF was performed by the total strain-controlled tests at 760 and 980 °C at different strain amplitudes to understand the effect of temperature and strain amplitude.

2 Experimental Procedures The nominal chemical composition of the DD11 was listed in Table 1. The cylinder specimen with [0 0 1] orientation was prepared by the withdrawal rate of 3 mm s−1 with a diameter of 15 mm and a length of 170 mm. The alloy was heat treatment at 1315 °C for 6 h and the aging 1130 °C for 4 h and 870 °C for 24 h. The specimens had a diameter of 6 mm and length of 90 mm which was used in the LCF tests. The orientation of the test bar deviation was within 8° by the X-ray backscattering Laue method. The tests were conducted by the MTS (Material Test System) servohydraulic testing machine under a total strain-range-control mode in air. The strain ratio was −1 and the strain amplitude (ε/2) was 0.6, and 1.0% which was measured and controlled by extensometer at 760 and 980 °C, and the strain rate of all of tests was 5 × 10−3 s−1 . After fatigue test, the longitudinal metallographic sections were observed by an FEI-nano 450 scanning electron microscope and the microstructural investigation was performed using an SEM (JEOL JXA-8100). The deformation microstructure after fatigue failure was examined by transmission electron microscopy (TEM) samples taken from transverse and longitudinal section of specimens by conventional grinding and twin-jet in a solution of 90% C2 H5 OH + 10% HClO4 at a temperature of −20 °C and a voltage of −20 mV using a JEM-2100F transmission electron microscope operating at a voltage of 200 kV.

Table 1 Chemical composition (wt%) of the DD11 Element

Cr

Co

W

Mo

Ta

Re

Al

Hf

B

C

Content

4

8

7

2

7

3

6

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0.1

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Fig. 1 The microstructure of DD11: a dendrite before heat treatment, b interdendritic and c γ and γ  phase after heat treatment

3 Results and Discussion 3.1 Microstructure The microstructure of DD11 was as shown in Fig. 1. The Fig. 1a and b were the dendrite and interdendritic morphology of the cast microstructure. There were plenty of eutectic at the interdendritic. The morphology of γ  was irregularity. The Fig. 1c were the microstructure of DD11 after heat treatment. The eutectic was disappeared and the morphology of γ  was cubic. The volume of fraction and scale of γ  phase were 65–70% and 380–425 nm, respectively.

3.2 Fatigue Life The LCF life of the DD11 alloy at the strain amplitude of 0.6 and 1.0% was 8972 and 374, respectively, under the condition of 980 °C. The LCF life was 58,180 and 109 at the strain amplitude of 0.6 and 1.0% at 760 °C. With increasing the total strain amplitude, the LCF life obviously decreased. However, the fatigue life of DD11 was not decreased as increasing the temperature at 0.6 and 1.0%, the fatigue life was lager at 760 than 980 °C when the strain amplitude was 0.6%, the fatigue life was lower at 760 than 980 °C at the strain amplitude of 1.0%.

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3.3 Cyclic Stress Response Behaviours The cyclic stress response behaviours at the strain amplitude of 0.6 and 1.0% at 760 and 980 °C were shown in Fig. 2. The cyclic softening and hardening behaviour could be observed which was closely related to the strain amplitude and temperature. When the temperature was 760 °C, the alloy behaved stable at 0.6%, as the strain amplitude increased to 1.0%, the cyclic stress response behaviour was cyclic hardening. At the strain amplitude of 0.6% at 980 °C, the DD11 alloy showed cyclic softening and the peak value of stress decreased about 100 MPa from the first cycle to the last stage. As the strain amplitude increased to 1.0%, the degree of softening decreased.

3.4 Fatigue Fracture Morphologies At 760 °C, the macroscopic fracture surface was inclined to the loading direction as shown in Fig. 3a and b. The macroscopic fracture surface was perpendicular to the loading direction under the strain amplitude of 0.6% at the 980 °C (Fig. 3c). The fatigue sources appeared in the margin of the fracture surface and were connected by sliding steps. There was the obvious boundary between the propagation region

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Fig. 3 The characteristics of fatigue surface under different loading conditions. a ε/2  0.6%, Nf  58,180, T  760 °C, b ε/2  1.0%, Nf  109, T  760 °C, c ε/2  0.6%, Nf  8972, T  980 °C, d ε/2  1.0%, Nf  374, T  980 °C

and the instant region at the strain amplitude of 1.0% at 980 °C (Fig. 3d), and the direction of crack propagation was along plane. The fracture surface showed that the sites of oxidation spalling which becomes the fatigue sources and finally result in fatigue failure at 980 °C (Fig. 4a, b, c). The orientation of the fatigue cracks was perpendicular to the fatigue striation at 760 and 980 °C as shown in Fig. 4d.

3.5 Deformed Microstructures After LCF The evolution of the γ /γ  coherent microstructure could reflect the deformation feature, thus the LCF tested specimens were sectioned longitudinally and the deformed microstructure near the fatigued specimens fracture surfaces was shown in Fig. 5. The direction of slip bands was oblique to the loading direction near the fracture surface. The distribution of slip bands was inhomogeneous. The cubic of γ  phase decreased at the strain amplitude of 0.6% at 980 °C.

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Fig. 4 The characteristics of fatigue surface under different loading conditions. a, b ε/2  0.6%, Nf  8972, T  980 °C, c ε/2  1.0%, Nf  374, T  980°C, d ε/2  1.0%, Nf  109, T  760 °C

The dislocation morphologies of the fatigue samples at the strain amplitude of 0.6% at 980 °C, as shown in Fig. 6a and b. In Fig. 6a, the dislocation of climb only appeared in the horizontal γ channel. In addition, the Fig. 6b indicated that there was cross-slip mode of the dislocation movement. As the strain amplitude of 1.0%, the dislocation movement mode was cross-slip which was benefit for the dislocation movement (as shown in Fig. 6c). As temperature decreased to 760 °C, the dislocation of cross-slip at 0.6% appeared in the γ channel and its density was lower than 980 °C (as shown in Fig. 6d). The dislocations cut into the γ  phase at the strain amplitude of 1.0% at 760 °C (as shown in Fig. 6e). At high temperature of 980 °C, the cyclic stress behaved cyclic softening and was related to the transformation of the morphology of γ  and the movement mode of dislocation. The cubic degree of γ  decreased under the cyclic loading condition at high temperature. The movement of dislocation by cross-slip and climb manners in the γ channel was easily and responsible for the softening. At the temperature of 760 °C, the dislocations which cut through into the γ  phase were responsible for the hardening.

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4 Conclusion In this study, the LCF of DD11 with different temperature and strain amplitude was carried out. Experimental results show that with the increase of strain amplitude the fatigue life decreases. However, at the strain amplitude of 0.6%, the fatigue life at 980 °C is lower than that at 760 °C, while at the strain amplitude of 1.0% the fatigue life at 760 °C is lower than that at 980 °C. At 980 °C and the strain amplitude of 0.6%, the cubic degree of γ  phase decreases which benefits for the movement of dislocations. The fracture surface is perpendicular to the loading direction. At 760 °C, the dislocation which cut through the γ  phase results in the cyclic hardening. In this case, the fracture surface inclines to the loading direction. Acknowledgements This study was financially supported by Beijing Natural Science Foundation (2184132).

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(a)

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Fig. 6 Dislocation configurations of the DD11 after LCF failure under different strain amplitudes. a, b ε/2  0.6%, Nf  58,180, T  980 °C, c ε/2  1.0%, Nf  109, T  980 °C, d ε/2  0.6%, Nf  8972, T  760 °C, e ε/2  1.0%, Nf  158

References 1. R.C. Reed, Superalloys: Fundamentals and Applications (Cambridge University Press, 2006) 2. M. McLean, Directionally Solidified Materials for High Temperature Service (The Metals Society, London, 1983) 3. K.H. David E. Laughlin, Physical Meallurgy, vol. III, 5th edn. (Elsevier, USA, 2014) 4. T.M. Pollock, Alloy design for aircraft engines. Nat. Mater. 15(8), 809–815 (2016) 5. H.U. Hong, J.G. Kang, B.G. Choi, I.S. Kim, Y.S. Yoo, C.Y. Jo, A comparative study on thermomechanical and low cycle fatigue failures of a single crystal nickel-based superalloy. Int. J. Fatigue 33(12), 1592–1599 (2011) 6. E. Silveira, G. Atxaga, A.M. Irisarri, Failure analysis of two sets of aircraft blades. Eng. Fail. Anal. 17(3), 641–647 (2010) 7. S. Barella, M. Boniardi, S. Cincera, P. Pellin, X. Degive, S. Gijbels, Failure analysis of a third stage gas turbine blade. Eng. Fail. Anal. 18(1), 386–393 (2011) 8. K. Makhlouf, J.W. Jones, Effects of temperature and frequency on fatigue crack growth in 18% Cr ferritic stainless steel. Int. J. Fatigue 15(3), 163–171 (1993) 9. A. Karabela, L.G. Zhao, J. Tong, N.J. Simms, J.R. Nicholls, M.C. Hardy, Effects of cyclic stress and temperature on oxidation damage of a nickel-based superalloy. Mater. Sci. Eng., A 528(19–20), 6194–6202 (2011) 10. M.R. Bache, W.J. Evans, M.C. Hardy, The effects of environment and loading waveform on fatigue crack growth in Inconel 718, Int. J. Fatigue 21(Supplement 1(0)),S69–S77 (1999) 11. N.T.R. Ohtani, M. Shibata, S. Taniyama, High temperature fatigue of the nickel-base singlecrystal superalloy CMSX-10. Fatigue Fract. Eng. Mater. Struct. 867–876 (2001) 12. Q.Z. Lizi He, X. Sun, H. Guan, H. Zhuangqi, K. Tieu, C. Lu, H. Zhu, Effect of temperature and strain amplitude on dislocation structure of M963 superalloy during high-temperature low cycle fatigue. Mater. Trans. 47(1), 67–71 (2016)

Oxidation Mechanism of Nb–Si-Based Ultra-High Temperature Materials Fengwei Guo, Yongwang Kang, Chenbo Xiao, Ming Li and Meiling Wu

Abstract The oxidation behavior of Nb–Si-based alloy (Nb–20Ti–16Si–3Cr–3Al–2Hf, at%) at 1250 °C in air was investigated. Severe selective oxidation occurred in the alloy, which mainly took place in Nb solid solution and accompanied with large amount of precipitations of Ti-rich oxide. External oxidation of Nb solid solution was the selective oxidation, composed of two steps of solid-phase reaction among the oxides. The morphology of oxide-scale layers varied since the solid-phase reaction between the oxides, and it became porous and multilayer finally. TiNb2 O7 , Ti2 Nb10 O29 , SiO2 , and TiO2 were the main phase constitutions of the oxide layer which were rich in Ti2 Nb10 O29 in inner oxide-scale layer and TiNb2 O7 in outer oxide-scale layer. The protection efficiency of the oxide-scale layer for the Nb–Si-based alloy at high temperature was weak. The weight gain during the beginning oxidation stage fitted the linear oxidation law with the rate of about 6 mg/(cm2 h).

1 Introduction Ultra-high temperature structural materials have been extensively studied to meet the urgent need of future aeronautical engine, which is expected to be more powerful and economic [1, 2]. Compare to other materials, Niobium–silicide-based alloy (Nb–Si alloy) has become the most promising alternative to Ni-based superalloy for its relatively low density, high strength at elevated temperature and moderate ductility [2, 3]. However, poor oxidation resistance becomes one of the main challenges of Nb–Si-based alloys [4–6]. Alloying has been proved as the most effective approach to improve the oxidation resistance of Nb–Si-based alloy [7–9], the elements usually used for alloying are Ti, Al, Cr, and Hf. Ti addition effectively suppresses severe oxidation at moderate temperature known as “pesting”, and Al, Cr, and Hf addition decreases oxygen solution in the alloy [8–10], but these elements still are not able F. Guo (B) · Y. Kang · C. Xiao · M. Li · M. Wu Science & Technology on Advanced High Temperature Structural Materials Laboratory, AECC Beijing Institute of Aeronautical Materials, 100095 Beijing, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_40

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to stop oxygen penetrating in the material, nor promote formation of continuous protective oxide layer above 1200 °C. This work aims to reveal the behavior of high-temperature oxidation of multielement Nb–Si-based alloy Nb-20Ti-16Si-3Cr-3Al-2Hf (at%). Weight-gain plots, oxides-phase diffraction and morphology of oxidation layer are obtained to analyze oxidation kinetics, two steps solid-phase reactions are raised to discuss the evolution of oxide layer.

2 Experimental Procedures Hemisphere ingots with nominal composition Nb-20Ti-16Si-3Cr-3Al-2Hf were arc melted using non-consumed tungsten electrode under 99.99% pure Ar atmosphere. Nine specimens named C1–C9 were electric discharging cut into size 7.5 × 6 × 3 mm from the ingots, and then been polished to 2000 meshes using abrasive paper. The static oxidation was examined at 1250 °C in air for 2.7, 6.3, 10, 18, 35 h, respectively, using alumina tube furnace, and the weight gain was measured by electric scales after specimen cool down to room temperature. Morphology of oxidized specimens was observed using scanning electron microscope (SEM). The oxide scale of the specimens after 35 h oxidation was milled then identified using X-ray diffractometer.

3 Results and Discussion The phase composition of the as-cast alloy is Niobium-based solid solution phase (Nbss ) and Niobium-silicide phase (Nb5 Si3 ). Nb5 Si3 phase distributes in the continuous Nbss phase as eutectic structure with few Nb5 Si3 primary phase.

3.1 Oxidation Rate The kinetics of oxidation (Fig. 1) is demonstrated by weight gain W (mg/cm2√ ) of unit area to time (h), as the weight gain W (mg/cm2 ) to root square of time ( h) is not linear that means the kinetics of oxidation does not follow the parabola law. Nine specimens were tested in order to raise the reliability of data. Specimens C8 shows much higher weight gain rate owning to containing lots of voids which are found by optical microscope and confirmed as casting defects, however, other eight plots show no significant difference with each other, indicates casting flaws to have little influence on oxidation in these eight specimens. Oxidation for 35 h consumed over 60% volume of base material that causes the plots to deviate significantly from linear line. The linear regression equation of the average weight gain W (mg/cm2 ) of unit area to time (h) is W  6t − 4.2 excluding the of

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Fig. 1 The weight gain of Nb–Si base alloy during oxidation process at 1250 °C

specimen C8 and the data at 35 h, and the confidence coefficient of regression is higher than 98%, which means that the oxidation rate is 6 mg/(cm2 h) at the beginning oxidation stage. The linear regression of weight gain indicates that oxidation of the Nb–Si-based alloy at 1250 °C fits linear law well, and the oxidation resistance is much weaker than Ni-based superalloy.

3.2 Selective Oxidation It has been widely proved that elements such as Ti, Cr, Al, and Hf are able to improve oxidation resistance of Nb-based alloy, however, the oxidation mechanism of metal solid solution and metalloid composites wide different to single-phase metallic alloy or metallic composites, as that although these elements solute in Nbss solid solution phase and silicide phase together, these two phases hardly oxide simultaneously at high temperature. Figure 2a demonstrates the backscattered electron image of material’s selective oxidation zone, region C is confirmed as external oxidation zone. The oxygen penetrates deeply into the base material (region A) through selective oxidation region B. Ti and Hf are oxidized and precipitate on Nb5 Si3 phase (Fig. 2b) because the interface of two phases could provide lower energy barrier for nucleation. EDS results demonstrate that oxygen solutes only in Nbss phase and the composition of Nb5 Si3 phase remain unchanged during the selective oxidation, which indicates that

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(d)

Fig. 2 Backscattered electron image of selective oxidation zone: a overview of selective oxidation zone; b preliminary oxidation zone in region A; c selective oxidation zone in region B; d the interface zone of inner layer oxide and outer layer oxide

the Ti and Hf oxides precipitate from Nbss phase and Nb5 Si3 phase did not participate in selective oxidation. Large amount of Ti oxides precipitate on the front surface of Nb5 Si3 phase along the direction of oxygen penetrating, and little elements of Ti, Al, and Cr remained in Nbss phase (Fig. 2c), in addition, Nb5 Si3 phase remains unchanged in region B. However, not all Nb5 Si3 phase is not oxidized in selective oxidation zone such as Ti-rich Nb5 Si3 phase shown in Fig. 3, which indicates that sufficient Ti addition should enhance the activity of reaction with oxygen. This conclusion is close to S. Mathieu’s research [6], even though the silicide in S. Mathieu’s research is Hf-rich (higher than 9 at%). After the alloying elements precipitate as oxides completely, the outer layer oxidation of Nbss phase and Nb5 Si3 phase begins as shown in Fig. 2d. The outer layer oxidation of Nbss phase and Nb5 Si3 phase performed at the same time since the interface of the outer oxidation layer and selective oxidation layer is generally straight.

Oxidation Mechanism of Nb–Si-Based Ultra-High Temperature …

(a)

403

(b)

Fig. 3 Ti-rich Nb5 Si3 phase oxidized in selective oxidation zone

3.3 External Oxidation Selective oxidation consumes alloying elements without changing the phase and structure of base alloy until little alloying elements remain in Nbss phase, then external oxidation takes place and starts to consume principle elements Nb and Si. However, the external oxidation of Nb–Si-based alloy is a long process because the oxygen beneath the oxide scale is far from enough to completely oxidize the alloy. The difference of solubility and reactivity of oxygen between two phases also decides the complexity of this process. The oxide scale of the specimens was peeled off and grinded into powder for Xray diffraction investigation as shown in Fig. 4. The result shows that the constitution of the oxide scale is mainly TiNb2 O7 , TiNb10 O29 , TiO2 , and SiO2 .

3.3.1

Oxidation of Nbss and Nb5 Si3 Phases

The external oxidation zone (Fig. 5a) is darker than the selective oxidation zone under backscattered electron imaging since the amount of oxygen raises significantly, and the metallic bonds of Nbss phase and covalent bonds of Nb5 Si3 phase are broken by oxygen shown in Fig. 5b. The composition of Region 1–6 in Fig. 5b was measured using EDS to monitor the migration of alloying elements, and the result is list in Table 1. The content of oxygen raises intensely in area 2 compared to unoxidized Region 1, and other element’s relative content remains unchanged, which indicates that the mechanism of oxidation is relying on the inward-diffusion of oxygen and the elements, Ti, Si, Al, Hf bound by covalent bonds diffuse outward against oxygen hardly. Figure 5c shows the process of Nb5 Si3 phase’s corrosion by oxygen. The interface between oxidized and unoxidized parts is clear, besides Nb5 Si3 phase separated into lots of extreme fine “needle-like” phases might be SiO2 (black needles) and Nb-based alloy/oxide (grey needles). The oxygen content for SiO2 in region 2

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TiNb2O7 Ti2Nb10O29 SiO2 TiO2

Fig. 4 X-ray diffraction patterns obtained for the analysis of oxide scale of Nb–Si alloy Table 1 EDS analyses performed on the Nb–Si-based alloy oxidized for 35 h at 1250 °C

Region

Chemical composition (at%) Nb

Ti

1(Nb5 Si3 )

47

16

2(Nb5 Si3 )

25

5

O 2.5 49

Si

Cr

Al

Hf

32

0

0

2.5

18

0

0

1

3(Nbss )

40

11

47

0

0

1.5

0.5

4(Nbss )

38

14

46

0

0

1.5

0.5

5(Nbss )

39

14

44

0

0

2

1

6(Nbss )

44

8

46

0

0

1

1

is much higher than that for oxides of Ti and Hf. These “needle-like” phases grow up quickly probably for reducing interfacial energy. The oxidation of Nb5 Si3 phase could be given in Eq. 1. Nb5 Si3 + 3O2  3SiO2 + 5Nb

(1)

Nbss phase oxidized at the same time with Nb5 Si3 phase shown in Fig. 5d, which formed Nb-oxide based on the reaction of Eq. 2. It is a nominal formula; however, the stoichiometric ratio varies in a large range actually. 4Nb + 5O2  2Nb2 O5

(2)

The synchronization of external oxidizing indicates that the activation of oxidation of Nbss phase in which most of alloying elements has diffused out is close to that of Nb5 Si3 phase at 1250 °C. This conclusion is hard to be proved because the process is kind of “black box” problem. The oxides of Nb5 Si3 phase are too small to investigate

Oxidation Mechanism of Nb–Si-Based Ultra-High Temperature …

(a)

(b) External/Selective

405

5

reaction interface

6

oxidation

3 1 4 2

crack

(c)

(d)

Ti oxides

Fig. 5 Backscattered electron image of the front of external oxidation

by EDS analysis in Fig. 5. Thermodynamic calculation is applied to explain the rationality in this work. Thermodynamic data of pure matter at 1500 K (close to 1250 °C) under standard state listed in Table 2 was used to calculate the free energy of Eqs. 1 and 2. The results show that the free energy change of both reactions is negative, which means that both Nb and Nb5 Si3 were oxidized under the standard state. However, the proportion of oxygen beneath the outer oxide scale is lower than standard state and there must be a bound where the matter is in equilibrium of oxidized and unoxidized state. The equilibrium oxygen partial pressure (POe 2 ) is the bound of oxygen content above, where Eqs. 1, 2 proceed forward and reverses backward. The formula of POe 2 is in:  POe 2  exp

G  RT

 (3)

where G  is standard Gibbs free energy which is independent of oxygen content. The result is 3.183 × 10−18 Pa for Eq. 1 and 4.898 × 10−17 Pa for Eq. 2. Although the POe 2 of Eq. 2 is 10 times larger than Eq. 1, these two values could be regarded

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Table 2 Thermodynamic data of pure matter at 1500 K under standard state

G 1500 K (kJ/mol) O2 Nb Nb5 Si3 Ti

−346.503 −86.619 −1132.145 −81.665

G 1500 K (kJ/mol) Si

−56.979

SiO2

−1047.869

Nb2 O5

−2295.6

TiO2

−1102.115

close for the simplification of thermodynamic calculation. One order’s difference of POe 2 might reflect in tiny distance between these two reacting interfaces shown in Fig. 5b, however, it is not able to be proved as the lack of diffusion data. Another explanation is that the reduction of POe 2 of Eq. 2 is caused by the great solid solution of oxygen in alloying element such as Ti and Hf in Nb5 Si3 phase. In addition, it is the reason for the selective oxidation of Ti-rich and Hf-rich Nb5 Si3 phase, when the oxygen content is much lower than the POe 2 for the oxidation of the normal Nb5 Si3 phase in Fig. 3.

3.3.2

Transformation of Oxides

The black particles (Ti-oxide) distribute in Nbss phase formed in selective oxidation vanishes in the Nb-oxide formed from Nbss phase shown in Fig. 5d. Ti content of Region 3 and 4 in Fig. 5b indicates that Ti-oxide merges together with Nb-oxide. The contrast difference between Region 3 and 4 in Nb-oxide indicates the different content of Ti. The reaction of the merge of Nb-oxide and Ti can be given in Eq. 4: 5Nb2 O5 + 2TiO2 → Ti2 Nb10 O29

(4)

The existence of Ti2 Nb10 O29 is proved by X-ray diffraction investigation in Fig. 4. In addition, the ternary oxide TiNb2 O7 is detected in outer oxide scale. The formation of TiNb2 O7 should follow behind the formation of Ti2 Nb10 O29 as TiNb2 O7 is main component of oxide scale, which means TiNb2 O7 is more stable after sufficient oxidation. Further X-ray diffraction investigation is performed separately on inner surface and outer surface of complete oxide scale which is thicker than 1 mm peeled off. The results are shown in Fig. 6. In the inner surface, the main component is Ti2 Nb10 O29 , with, TiO2 and HfO2 detected. The patterns of TiO2 and HfO2 disappeared in the outer surface which means that the oxide has been transformed to TiNb2 O7 completely. The result of X-ray diffraction analysis shows that Ti2 Nb10 O29 reacts with TiO2 and HfO2 and formed TiNb2 O7 after sufficient oxidation. The formation of TiNb2 O7 is given in Eq. 5: Ti2 Nb10 O29 + 3TiO2 → 5TiNb2 O7

(5)

Oxidation Mechanism of Nb–Si-Based Ultra-High Temperature …

Inner surface

TiNb2O7 HfO2

407

Ti2Nb10O29 TiO2

Outer surface

Fig. 6 X-ray diffraction patterns of inner face and outer face of oxide scale

The transformation of oxides does not only change the composition but also the morphology of the oxide layer. The oxidation of Nb5 Si3 phase produced large amount of small oxide particles (Eq. 1) and the structure of the product remained dense in Fig. 7a, e. However, the reaction of oxides from Nbss phase forms lots of pores that make the layer more porous (Fig. 7b, d, f). It takes a long time for Nb5 Si3 phase to oxidize completely and leaves a thick reaction layer C in Fig. 7e. When the oxidation of Nb5 Si3 phase is finish, all oxides produced by Nbss and Nb5 Si3 phase start reacting to form needle-like TiNb2 O7 grains (Eq. 5) in dense array perpendicular to the surface. The SiO2 particles formed by Nb5 Si3 phase distribute among the gaps of TiNb2 O7 grains that induces the layer dense. The driving force of the transformation in Eq. 5 remains unknown and it leaves lots of pores along the boundary in Fig. 7e, f that caused the oxide layer easily spalls during high-temperature oxidation.

4 Summary The oxidation mechanism of multielement alloying Nb–Si-based alloy (Nb-20Ti16Si-3Cr-3Al-2Hf, at%) under 1250 °C in air was investigated including the kinetic behavior, the morphology of selective oxidation and external oxidation, the oxides composition and the reactions among oxides. The main conclusions obtained in this study are listed as follows: (1) Severe selective oxidation takes place in all as-cast alloys, where the lots of Ti-oxide form on Nb solid solution. The Nb5 Si3 phase remains unchanged in selective oxidation layer, at the same time Ti-rich Nb5 Si3 phase participates

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(a)

(b)

(c)

(d)

(e)

(f)

Fig. 7 Morphology of oxide layers of two steps solid-phase reaction: a, c, e backscattered electron image; b, d, f secondary electron image

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409

during selective oxidation. Ti retards the diffusion of oxygen to Nb–Si-based alloy and increases the reaction activity of Nb5 Si3 with oxygen. (2) The weight gain of this alloy follows linear law with regression equation W  6t − 4.2 at the beginning oxidation stage. The oxide products are TiNb2 O7 , Ti2 Nb10 O29 , TiO2 , and SiO2 . The external oxidation of Nb solid solution and Nb5 Si3 phase take place separately and almost at the same time. The external oxidation of Nb solid solution accompanied by the reaction of oxide product: 5Nb2 O5 + 2TiO2 → Ti2 Nb10 O29 , while the external oxidation of Nb5 Si3 phase takes a long time to complete. The oxides produced by Nb solid solution and Nb5 Si3 phase reacting with O form a dense layer of needle-like TiNb2 O7 grains, however, they leave lots of pores at the boundary which should lead to the easy oxide-scale spallation during high-temperature oxidation. (3) The oxide layer of the Nb–Si-based multielement alloy formed at 1250 °C is a multilayer with porous, which is not able to prevent the oxygen to penetrate into Nb–Si substrate. Acknowledgements This work was supported by National Key R&D Program of China (No. 2017YFB0702904) and National Natural Science Foundation of China (No. 51601183).

References 1. B.P. Bewlay, M.R. Jackson, J.C. Zhao et al., A review of very-high-temperature Nb-silicidebased composites. Metall. Mater. Trans. A 34A(10), 2043–2052 (2003) 2. C. Roger, Reed, The Superalloys Fundamentals and Applications (Cambridge University Press, 2006) 3. T. Murakami, S. Sasaki, K. Ichikawa et al., Microstructure, mechanical properties and oxidation behavior of Nb-Si-Al and Nb-Si-N powder compacts prepared by spark plasma sintering. Intermetallics 9(7), 621–627 (2001) 4. M.D. Moricca, S.K. Varma, High temperature oxidation characteristics of Nb-10 W-XCr alloys. J. Alloy. Compd. 489(1), 195–201 (2010) 5. B.I. Portillo, S.K. Varma, Oxidation behavior of Nb-20Mo-15Si-5B-20Ti alloy in air from 700 to 1300 °C. J. Alloy. Compd. 497(1–2), 68–73 (2010) 6. S. Mathieu, S. Knittel, P. Berthod, S. Mathieu, M. Vilasi, On the oxidation mechanism of niobium-base in situ composites. Corros. Sci. 60(12) (2012) 7. J. Geng, P. Tsakiropoulos, G.S. Shao, Oxidation of Nb-Si-Cr-Al in situ composites with Mo, Ti and Hf additions. Mater. Sci. Eng. A. Struct. Mater. 441(1–2), 26–38 (2006) 8. A. Mueller, G. Wang, R.A. Rapp et al., Deposition and cyclic oxidation behavior of a protective (Mo, W)(Si, Ge) 2 coating on Nb-base alloys. J. Electrochem. Soc. 139(5), 1266–1275 (1992) 9. K. Zelenitsas, P. Tsakiropoulos, Study of the role of Al and Cr additions in the microstructure of Nb-Ti-Si in situ composites. Intermetallics 13(10), 1079–1095 (2005) 10. W.Y. Kim, I.D. Yeo, M.S. Kim et al., Effect of Cr addition on microstructure and mechanical properties in Nb-Si-Mo base multiphase alloys. Mater. Trans. 43(12), 3254–3261 (2002)

Microstructure and Mechanical Properties of Aluminum Alloy with Ultra-high Strength Prepared by Spray Forming Shuhui Huang, Baiqing Xiong, Yong’an Zhang, Zhihui Li, Xiwu Li, Hongwei Liu, Hongwei Yan, Lizhen Yan and Kai Wen Abstract In this paper, Al–11Zn–3Mg–2Cu–0.2Zr alloy was prepared by sprayforming, and then it was treated by defect reduction, hot worked and heat treatment. Microstructure and mechanical properties of the aluminum alloy were researched during the whole processing by metallographic microscope, scanning electron microscope (SEM), and tensile test. Average grain size of spray-formed aluminum alloy was about 10–20 µm due to the high cooling rate closing 1000 K/s, and shrinkage porosity and cavity was inevitable in spray-formed ingot because of nitrogen as carrier. Through hot isostatic pressing and hot extrusion, the relative density increased from 87– 90% to almost 99–100%, but it decreased again to about 98% after solution, which meant some porous defects could not be eliminated thoroughly. Comparing with casting, the original grain of spray-formed ingot was much smaller, which was unsuited for plastic processing with large deformation. A large plastic deformation caused lots of recrystallization, which weakened the mechanical properties. In the areas without defects, the best properties of the spray-formed alloy were 825 MPa in tensile strength, 808 MPa in yield strength, and 9% in elongation.

1 Introduction 7xxx series aluminum alloy is one of the highest strength aluminum alloys, and the strength of 7xxx series aluminum alloy will continue to raise with content of the main alloying elements (Zn, Mg, Cu, and so on) increase. The cooling rate of semicontinuous casting process is about 10 K/s, which restricts content of the main alloying elements increasing because of insufficient undercooling. Large precipitate will form and cannot be restored into matrix by heat treatment, which is harmful to mechanical properties of aluminum alloy. So when the 7xxx series aluminum alloy is made by traditional casting process, the content of the main alloying elements cannot

S. Huang (B) · B. Xiong · Y. Zhang · Z. Li · X. Li · H. Liu · H. Yan · L. Yan · K. Wen State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co., Ltd., Beijing 101407, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_41

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exceed 15 wt% normally, which leads to the strength of 7xxx series aluminum alloy hard to exceed 650 MPa [1–3]. In 1968, Singer of Swansea University invented a kind of rapid solidification technology called spray forming [4–6]. The cooling rate of spray forming process is about 103 –104 K/s, so the content of the main alloying elements can be increased to more than 15 wt%, and the strength will be up to 700 MPa or higher. Compared with casting ingot, the aluminum alloy ingot prepared by spray forming has the following microscopic features, higher alloying elements solid solubility, finer grain, more uniform composition, and organization [7–11]. Of course, spray forming also has its inherent problems, and there are inevitable pore defects in the ingot [11–13]. The purpose of this work is to prepare an aluminum alloy with tensile strength exceeding 800 MPa by spray forming, and provide reference for the application of spray forming process in the field of ultra-high strength aluminum alloy preparation and processing.

2 Materials and Methods In this paper, the experimental process of the new Al–11Zn–3Mg–2Cu–0.2Zr alloy is spray forming, defect reduction, hot-worked, heat treatment and performance test. The spray forming process is carried out on a self-made machine. Microstructures are observed by Carl Zeiss Zxiovet 200MAT optical microscopy (OM) and JSM-7001 field emission scanning electron microscope (SEM). A NETZSCH STA 409C/CD equipment is used to get differential scanning calorimetry (DSC) curve in the temperature range 25–600 °C in an Argon protection with heating rate of 0.2 °C/s. 1250 t horizontal extruder is used for extrusion experiments at about 400 °C with different ratios.

3 Results and Discussion 3.1 Microstructure Evolution During Hot-Worked Spray-formed ingots are shown in Fig. 1, which are all about 200 × 300. The original microstructure of spray-formed ingot is shown in Fig. 2, as can be seen that the grain size is about 10–20 um and lots of 10–100 µm porous defect distribute diffusely. The SEM photo of spray-formed ingots after 450 °C/8 h hot isostatic pressing is shown in Fig. 3. Compared with Fig. 2, most of the porous defect is eliminated through hot isostatic pressing, but some residues can still be observed. Endothermic peak of the alloy by spray forming appears at 473.5 °C in differential scanning calorimeter curve, as shown in Fig. 4. Homogenization is carried

Microstructure and Mechanical Properties …

Fig. 1 Spray-formed ingots

Fig. 2 The original microstructure of spray-formed ingot

Fig. 3 SEM photo of spray-formed ingots after 450 °C/8 h hot isostatic pressing

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Fig. 4 DSC curve of spray-formed alloy

473ºC/36h

473ºC/48h

Fig. 5 The SEM photo of spray-formed ingots after homogenization

out at 473 °C according to DSC, and the SEM photo of spray-formed ingots after homogenization is shown in Fig. 5. Precipitate phase is basically dissolved after 48 h. Extrusion is implemented at 400 °C with extrusion ratio of 15:1 and 30:1. Solution treatment is carried out at 473 °C/3 h, 4 h, and 5 h, then the EBSD photo is shown in Fig. 6. Substructure is obviously coarsened with the extension of solution time and the effect of extrusion ratio on microstructure is shown in Fig. 7. As can be seen that, when the extrusion ratio is 30:1, recrystallization causes coarse tissue. In sprayformed aluminum alloy, the grain size is about 10–20 um, which is much smaller than that of 100–200 um in ordinary semicontinuous casting because of larger degree of sub-cooling in spray forming. Fine grains have many advantages, but this causes recrystallization to occur more easily during deformation.

Microstructure and Mechanical Properties …

415

473ºC/2h

Fig. 6 The EBSD photo of the extruded alloy with ratio of 15:1 after solution

15:1

30:1

Fig. 7 Optical microstructure of the alloy with different ratios after 473 °C/2 h solution

3.2 Mechanical Properties After peak aging, the mechanical properties of spray-formed alloy were tested by tensile test. The mechanical properties of the material vary greatly at different locations. In areas without defects, the best performance of the spray-formed alloy is 825 MPa tensile strength, 808 MPa yield strength, and 9% elongation, respectively. But in areas with lots of defects, the worst performance of the alloy is less than 600 MPa tensile strength and 4% elongation, respectively. The SEM photo of tensile fracture is shown in Fig. 8, whose morphological features accords to intergranular fracture, and some 10–20 µm porous defect is found. Most of the porous defect can be eliminated after hot isostatic pressing and extrusion, but it cannot be completely eliminated. Some porous defect contains gas, and it expands again during heat treatment. Residual defect is harmful for the mechanical properties of spray-formed alloy obviously. Ultrasonic flaw detection cannot detect this kind of defect, because it is too small for ultrasonic wavelength and cannot be measured. Drainage method is adopted to measure the evolution of density and relative density of spray-formed alloy, and the result is shown in Table 1. The density of cast

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Fig. 8 Tensile fracture of spray formed 7000 aluminum alloy Table 1 The evolution of density and relative density of spray-formed alloy No.

Density (g/cm3)

Relative density (%)

O

HIP

HOM

EXT

SOL

O

HIP

HOM

EXT

SOL

1

2.52

2.88

2.84

2.88

2.84

87.2

99.6

98.3

99.5

98.2

2

2.60

2.87

2.83

2.87

2.84

89.9

99.3

97.9

99.3

98.4

3

2.59

2.88

2.84

2.86

2.85

89.6

99.5

98.1

98.9

98.6

Fig. 9 X-ray diffraction spectrum of secondary phases in spray-deposited alloy

alloy with the same chemical composition is used to measure relative density, which is 2.89 g/cm3 . The result shows that most of the porous defect can be eliminated after hot isostatic pressing and extrusion, but it expands again during heat treatment because of containing gas. Spray-forming is carried on in non-vacuum environment, and nitrogen is used for carrier. The original spray-formed ingot is processed into powder for X-Ray diffraction, which is shown in Fig. 9. Characteristic diffraction peak of Al2 O3 , MgO, AlN, and Mg3 N2 are founded, which means some oxides and nitrides exist in sprayformed and the mechanical properties of spray-formed alloy are limited by them.

Microstructure and Mechanical Properties …

417

4 Summary (1) Al–11Zn–3Mg–2Cu–0.2Zr is prepared by spray-forming, and the grain size is about 10–20 µm and lots of 10–100um porous defect distribute diffusely in ingot. (2) Most of the porous defect can be eliminated after hot isostatic pressing and extrusion, but it expands again during heat treatment because of containing gas. The same law is proved by the evolution of density and relative density of spray-formed alloy during processing. (3) In the areas of the alloy without defects, the best performances of the sprayformed alloy are 825 MPa in tensile strength, 808 MPa in yield strength, and 9% in elongation. (4) Residual defect, oxides and nitrides are all harmful for the mechanical properties of spray-formed alloy obviously.

References 1. B.Q. Xiong, Y.A. Zhang, B.H, Zhu. Research on ultra-high strength Al-11Zn-2.9Mg-1.7Cu alloy prepared by spray forming process. Mater. Sci. Forum V475–479, 2785–2788 (2005) 2. E.J. Lavernia, T.S. Srivatsan, The rapid solidification processing of materials: science, principles, technology, advances, and applications. Mater. Sci. 45, 287–325 (2010) 3. Y.A. Zhang, B.H. Zhu, B.Q. Xiong, Research on ultra-high strength Al-10.8Zn-2.9Mg-l.7Cu alloys from spray forming. Chin. J. Rare Metals 27, 609–613 (2003) 4. E.J. Lavernia, J.D. Ayers, T.S. Srivatsan, Rapid solidification processing with specific application to aluminium alloys. Int. Mater. Rev. 37, 1–44 (1992) 5. M. Gupta, J. Juarez-Islas, W.E. Frazier, Microstructure, excess solid solubility and elevated temperature mechanical behavior of spray-atomized and codeposited Al-Ti-SiCp. Metall. Mater. Trans. 23, 719–736 (1992) 6. T.S. Srivatsan, T.S. Sudarshan, E.J. Lavernia, Processing of discontinuously-reinforced metal matrix composites by rapid solidification. Prog. Mater Sci. 39, 317–409 (1995) 7. T.S. Sidhu, R.D. Agrawal, S. Prakash, Hot corrosion of some superalloys and role of highvelocity oxy-fuel coating-a review. Surf. Coat. Technol. 198, 441–446 (2005) 8. A.R.E. Singer, Principles of spray rolling of metals. Metals Mater. 4, 246 (1970) 9. E.J. Lavernia, E.M. Gutierrez, J. Szekely, Spray deposition of metals. Mater. Sci. Eng. A 98, 381–394 (1988) 10. E. Salamci, Spray casting. Gazi Univ. J. Sci. 17(2), 155–173 (2004) 11. P. Lensfeld, Microstructure and mechanical behavior of spray deposited Zn modified 7xxx Series Al alloys. Int. J. Rapid Solidification 8, 237–265 (1995) 12. P. Mathur, S. Annavarapu, D. Apelian, Spray casting: an integral model for process understanding and control. Mater. Sci. Eng. A 142, 261–276 (1991) 13. Q. Xu, E.J. Lavernia, Influence of nucleation and growth phenomena on microstructural evolution during droplet based deposition. Acta Mater. 49, 3849–3861 (2001)

Effect of Pre-aging Technology on Microstructure and Mechanical Properties of 6111 Aluminum Alloy Hongwei Liu, Shuhui Huang, Baiqing Xiong, Yong’an Zhang, Zhihui Li, Xiwu Li, Hongwei Yan, Lizhen Yan and Kai Wen

Abstract 6xxx series aluminum alloy is one of the ideal automotive lighting materials for its high strength, excellent formability, good corrosion resistance, and weldability. In this paper, the effect of pre-aging technology on microstructure and properties of 6111 aluminum alloy was investigated by using TEM, tensile, and Erichsen test. The results showed that better properties of 6111 alloy could be obtained through 140 °C/10 min pre-aging treatment within 30 min after solution treatment (named T4P treatment). The n, r, I E , and yield strength values of 6111-T4 alloy were 0.31, 0.62, 8.06 mm, and 149 MPa, respectively, while those of 6111-T4P alloy were 0.33, 0.76, 8.45 mm, and 133 MPa, respectively. After simulated paint baking at 170 °C/30 min, the yield strength of 6111-T4 and 6111-T4P alloys increased to 154 and 212 MPa, respectively. Compared with T4 treatment, the pre-aging treatment reduced precipitating temperature of β phase and promoted precipitation during simulated paint baking. Pre-aging treatment benefits press forming of automotive body sheet and enables strengthening of the materials after simulated paint baking.

1 Introduction 6xxx series alloys are increasingly used for automotive body panels because of their high strength-to-weight ratio, good formability, and corrosion resistance [1, 2]. They can be strengthened through artificial aging at elevated temperatures, resulting in the formation of nanosized, coherent, or semi-coherent metastable precipitates which act as obstacles to dislocation movement in the Al matrix [3–5]. Among the precipitates of the 6xxx alloys, the needle-like β precipitates are aligned along 100Al and the dominant strengthening phases [6]. These alloys for automobile industry applications would have low yield strength and good formability while high yield strength after paint baking process for in-service dent. Unfortunately, the relatively low temperature and short holding time of the typical paint baking H. Liu · S. Huang (B) · B. Xiong · Y. Zhang · Z. Li · X. Li · H. Yan · L. Yan · K. Wen State Key Laboratory of Nonferrous Metals and Processes, GRIMAT Engineering Institute Co., Ltd., Beijing 101407, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_42

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processes cannot fully exploit the age hardening potential of the Al–Mg–Si alloys. Besides, in practical manufacture, storing Al–Mg–Si alloys at room temperature prior to paint baking is usually unavoidable and thus natural aging (NA) would happen. This generates two problems for industrial application of Al–Mg–Si alloys. First, the formation of clusters during NA results in the enhanced strength, which can reduce formability during stamping [7]. Second, these clusters cause the so-called negative effect of NA on BH response [8, 9]. Hence, maximal suppression of NA and remarkable improvement in BH response are essential for successful application of 6xxx alloys in automotive industry. Nowadays, pre-aging treatment immediately after quenching has been developed to improve the mechanical properties of these alloys. It is known that clusters with a uniform Mg/Si are formed during PA and they can easily transform into β phases at paint baking temperature, leading to an enhanced BH response [10, 11]. The formation of clusters during PA results in the lower vacancy and solute concentrations, which prevent the formation of clusters during subsequent NA [9, 10]. As a consequence, PA has a mitigating influence on the negative effect of NA. In this work, the effect of different pre-aging treatment on microstructures and mechanical properties of 6111 alloy is studied. The purpose of this study is optimization of the suitable pre-aging treatment for 6111 alloy to ensure the good formability and enhanced bake hardening response.

2 Experimental Procedure A commercial 6111 alloy was studied in this work. The studied alloy was cast, homogenized, hot-rolled and cold rolled to 1 mm thick sheets. All samples were solution treated for 0.5 h at 550 °C and quenched into room temperature water. One set of samples was immediately pre-aged at different temperature for different hold times. The other set of samples was naturally aged for 2 weeks after water quenching. All samples were carried out the paint bake treatment (170 °C for 30 min). The tensile tests of the alloy were performed on the natural aged, pre-aged, and paintbaked alloy sheets. The strain hardening index (n), the plasticity strain ratio (r) and Erichsen value (I E ) were measured to characterize formability of the natural aged and pre-aged samples. Microstructure investigations were done by transmission electron microscope (TEM). All TEM images shown in this work were taken along 001Al directions. Differential scanning calorimetry (DSC) under an argon atmosphere is conducted at a heating rate of 10 °C/min. The TEM specimens were electro-polished with an electrolyte consisted of 1/3 HNO3 in methanol at a temperature of 30 °C.

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3 Results and Discussions 3.1 Tensile Results

360 345 330 315 300 285 270 255 240 225 210 195 180 165 150 135 120 105 90

(a)

5min 10min r 20min

T4

10 20 30

Rm Rp0.2 /MPa

Rm Rp0.2 /MPa

Strength changes of 6111 alloy before and after the paint bake treatment are shown in Fig. 1. As seen in Fig. 1a, the yield strength and ultimate tensile strength of alloy decrease slowly and then increase with the increasing pre-aging temperature. Preaging at 140 °C has the lowest strength before paint bake treatment. As shown in Fig. 1b, the yield strength and ultimate tensile strength after paint bake treatment increase with increasing pre-aging temperature from 100 to 140 °C while decreasing when the pre-aging temperature increases to 180 °C. In addition, 6111 alloy has the lower yield strength during pre-aging of different temperatures and holding times and higher yield strength after subsequent paint bake treatment, compared with T4 treatment. Those results indicate that pre-aging treatment can offer an enhanced bake hardening response without impairing the formability. In practical manufacture, storing 6xxx alloys at room temperature prior to paint baking is usually unavoidable and thus natural aging (NA) would happen. Serizawa et al. [11] have been reported that cluster(1) forms and strengthens the alloy during NA. By contrast, cluster(2) is formed during pre-aging, which can transform to β phase because of their similar Mg/Si ratio. This leads to an enhanced bake hardening response. On the other hand, the formation of cluster(2) during PA results in the lower vacancy and solute concentrations, which prevent the formation of clusters during subsequent NA [9, 10]. As a consequence, PA has a mitigating influence on the negative effect of NA. In this study, the higher bake hardening response of 6111 alloy caused by pre-aging treatment can be attributed to the formation of cluster(2). Figure 1 also shows that pre-aging at 140 °C has a minimum pre-aging strength and relatively higher pain-baked strength, suggesting that 140 °C is the ideal process to guarantee the formability and bake hardening response of 6111 alloy. The short pre-aging time (140 °C/5 min) cannot guarantee the stability of materials. The long

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Fig. 1 Change in yield strength and ultimate tensile strength of 6111 alloy before and after the paint bake treatment. a Natural aging and different pre-aging, b paint bake condition

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pre-aging time (140 °C/20 min) do not result in obvious improvement yield strength before and after pain bake. Thus, the reasonable pre-aging process in the present work is chosen as 140 °C/10 min.

3.2 Formability and Bake Hardening Response In Sect. 3.1, pre-aging at 140 °C for 10 min have a low yield strength and enhanced bake hardening response. In this section, the strain hardening index (n), the plasticity strain ratio (r), and Erichsen value (I E ) are used to characterize formability of 6111 alloy during pre-aging at 140 °C for 10 min. Pre-aging of 140 °C/10 min is named as T4P in this paper. The formability of 6111 alloy under T4 and T4P condition is shown in Fig. 2. The n value of alloy is 0.33, the r value is 0.76, and the I E value is 8.45 mm under T4P condition, which is higher than that of alloy under T4 condition (n value for 0.31, r value for 0.62, and I E value for 8.06 mm). This result indicates that pre-aging of 140 °C/10 min can improve the formability of 6111 alloy. Mechanical properties of 6111 alloy during different heat treatments are shown in Table 1. A stimulated paint bake treatment is referred to as artificial aging at 170 °C for 30 min, which is designed as PB in this work. Yield strength (YS) of T4 is 149 MPa and YS increases to 154 MPa after PB treatment, a 5 MPa of paint bake hardening response (PBR) is achieved. On the other hand, the YS under T4P condition are 133 MPa and YS changes to 212 MPa. PBR (79 MPa) of T4P is higher compared with T4.

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Fig. 2 Formability of 6111 alloy under T4 and T4P condition

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3.3 DSC Analysis The DSC curves of 6111 alloy immediately after quenching and T4 treatment are shown in Fig. 3. For as-quenched sample, the exothermic peaks occurred can be correlated with the various stages of the precipitation process: exothermic peak 1 (clusters), exothermic peak 2 (GP zones), exothermic peak 3 (β ), exothermic peak 4 (β ) [5, 12]. The enthalpy of β precipitate is 5.02 J/g for T4P sample. The endothermic peak between peak 2 and peak 3 is attributed to the dissolution of GP zones [5]. Interestingly, the exothermic peak 1 and 2 of T4 sample disappear and the enthalpy of β increases to 6.81 J/g. In addition, a significant endothermic peak between 125 and 200 °C is observed in T4 sample, which is caused by the dissolution of cluster(1) [13]. The cluster(1) formed after T4 cannot act as nuclei for the β phase. The increased enthalpy of β can be explained by the dissolution of cluster(1) on β phase [14]. The DSC curves of 6111 alloy of T4 and T4P samples are shown in Fig. 4. The T4P sample has only exothermic peak 3 (β ) and exothermic peak 4 (β ). It is noteworthy that the β precipitation peak of T4P sample shifts to lower temperature compared with T4 sample. The cluster(2) formed during pre-aging can transform easily into the β phase because of the compositional similarity between the cluster and β phase. The formation of cluster(2) greatly promotes the precipitation of β . This leads to lower β precipitation temperature.

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Fig. 3 DSC curves of 6111 alloy after quenching and T4

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3.4 TEM Observation TEM bright field images and (001) SAED patterns for the 6111 alloy after T4 and PB treatment are displayed in Fig. 5. There is no indication of precipitate in the bright field images, which is consistent with only diffraction spots of the Al matrix observed in the corresponding SAED patterns. The strength of 6111 alloy after T4 and PB treatment should be ascribed solely to some clusters. The strengthening due to clusters has been reported in 6xxx alloys [15, 16]. These clusters are difficult to resolve by TEM examinations because of their small size and presumed coherence with the Al matrix [17]. The clusters(1) formed during T4 treatment are not favorable nucleation sites for β phases and dissolving in paint baking condition. In addition, lower solute and vacancy concentrations in matrix inhibit nucleation of suitable precursors of β phase. The clusters, GP zones and β phases have different strengthening effects on 6xxx alloy. Considering unit precipitation particle, the strengthening of the clusters, GP zones and β phases on alloy can be describable to be: β phases > GP zones > solute clusters [18]. After T4 and PB treatment, β phases are not observed in 6111 alloy, leading to the low bake hardening response. Figure 6 shows the TEM and HRTEM images of 6111 alloy after T4P + PB. As shown in Fig. 6a, the needle-shaped and point-shaped precipitates are observed in the bright field images. The needle-like precipitates are aligned along the Al direction, the corresponding SAED patterns display streaks along the 001Al direction (Fig. 6b). Considering the shape, orientation relationship, SAED patterns [19], the needle-shaped precipitate are identified as β phase. The point-shaped precipitates (about 3–4 nm) should be the cross sections of the needle β precipitates which are aligned with the 001Al direction parallel to the electron beam. The point-shaped

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Fig. 5 TEM bright field images and (001) selected area electron diffraction (SAED) patterns obtained from 6111 alloy after T4 + PB treatment

precipitates with the size of 2 nm can be identified as GP zones. As shown in Fig. 6c, the length of the typical β precipitate is about 8–9 nm and the diameter of GP zones is about 2 nm. The enlarged images of the precipitates (Fig. 6d) suggest that the precipitates are fully coherent with the matrix. Cluster(2) with a uniform Mg/Si are formed during T4P and they can easily transform into β phases at paint baking temperature, leading to an enhanced BH response.

4 Conclusions The effect of pre-aging treatment on formability and mechanical properties of 6111 alloy was studied in this work. The following conclusions can be drawn: 1. The yield strength and ultimate tensile strength of alloy decrease slowly and then increase with the increasing pre-aging temperature. After subsequent paint bake treatment, the yield strength and ultimate tensile strength show an opposite trend with the increasing pre-aging temperature. 2. The reasonable pre-aging process in the present work is chosen as 140 °C/10 min (T4P). This T4P treatment can improve the formability of alloy and enhance bake hardening response. 3. The TEM result shows that after T4 + PB treatment, no precipitate is observed, but GP zones and β phase are formed. The DSC analysis indicates that T4P treatment prompts the formation of β phase and enhances the bake hardening response.

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Fig. 6 TEM and HRTEM images of 6111 alloy after T4P + PB. a TEM bright field image, b the corresponding SAED patterns, c HRTEM image, and d enlarged images of precipitates shown in c and corresponding FFT patterns

References 1. M.A. van Huis, J.H. Chen, M.H.F. Sluiter, H.W. Zandbergen, Phase stability and structural features of matrix-embedded hardening precipitates in Al-Mg-Si alloys in the early stages of evolution. Acta Mater. 55, 2183–2199 (2007) 2. S. Pogatscher, H. Antrekowitsch, H. Leitner, T. Ebner, P.J. Uggowitzer, Mechanisms controlling the artificial aging of Al-Mg-Si alloys. Acta Mater. 59, 3352–3363 (2011) 3. C.S. Tsao, C.Y. Chen, U.S. Jeng, T.Y. Kuo, Precipitation kinectics and transformation of metastable phases in Al-Mg-Si alloys. Acta Mater. 54, 4621–4631 (2006) 4. W.C. Yang, L.P. Huang, R.R. Zhang, M.P. Wang, Z. Li, Y.L. Jia, R.S. Lei, X.F. Sheng, Electron microscopy studies of the age-hardening behaviors in 6005A alloys and microstructural characterizations of precipitstes. J. Alloys Compd. 514, 220–233 (2012)

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5. G.A. Edwards, K. Stiller, G.L. Dunlop, M.J. Couper, The precipitation sequence in Al-Mg-Si alloys. Acta Mater. 46, 3893–3904 (1998) 6. W.F. Miao, D.E. Laughlin, Effects of Cu content and preaging on precipitation characteristics in aluminum alloy 6022. Metall. Mater. Trans. A 31, 361–371 (2000) 7. R. Prillhofer, G. Rank, J. Berneder, H. Antrekowitsch, P.J. Uggowitzer, S. Pogatscher, Property criteria for automotive Al-Mg-Si sheet alloys. Materials 7(7), 5047–5068 (2014) 8. L. Zhen, S.B. Kang, The effect of pre-aging on microstructure and tensile properties of AlMg-Si alloys. Scr. Mater 36, 1089–1094 (1997) 9. M. Torsæter, H.S. Hasting, W. Lefebvre, C.D. Marioara, J.C. Walmsley, S.J. Andersen, R. Holmestad, The influence of composition and natural aging on clustering during preaging in Al-Mg-Si alloys. J. Appl. Phy. 108, 0735527 (2010) 10. Y. Aruga, M. Kozuka, Y. Takaki, T. Sato, Evaluation of solute clusters associated with bakehardening response in isothermal aged Al-Mg-Si alloys using a three-dimensional atom probe. Metall. Mater. Trans. A 45, 5906–5913 (2014) 11. Serizawa, S. Hirosawa, T. Sato, Three-dimensional atom probe characterization of nanoclusters responsible for multistep aging behavior of an Al-Mg-Si alloy. Metall. Mater. Trans. A 39, 243–251 (2008) 12. L. Zhen, S.B. Kang, DSC analyses of the precipitation behavior of two Al-Mg-Si alloys naturally aged for different times. Mater. Lett. 37, 349–353 (1998) 13. L.P. Ding, Y. He, Z. Wen, P.Z. Zhao, Z.H. Jia, Q. Liu, Optimization of the pre-aging treatment for an AA6022 alloy at various temperatures and holding times. J. Alloys Compd. 647, 238–244 (2015) 14. V.N. Grau, A. Cuniberti, A. Tolley, V.C. Riglos, M. Stipcich, Solute clustering behavior between 293 K and 373 K in a 6082 aluminum alloy. J. Alloys Compd. 684, 481–487 (2016) 15. R.K.W. Marceau, A. de Vaucorbeil, G. Sha, S.P. Ringer, W.J. Poole, Atom probe tomography and yield stress modelling. Acta Mater. 61, 7285–7303 (2013) 16. M.J. Starink, L.F. Cao, P.A. Rometsch, A model for the thermodynamics of and strengthening due to co-clusters in Al-Mg-Si-based alloy. Acta Mater. 60, 4194–4207 (2012) 17. F.A. Martinsen, F.J.H. Ehlers, M. Torsæter, R. Homestad, Reversal of the negative natural aging effect in Al-Mg-Si alloys. Acta Mater. 60, 6091–6101 (2012) 18. H. Li, W. Liu, Nanoprecipitates and their strengthening behavior in Al-Mg-Si alloy during the aging process. Metall. Materi. Trans A 48, 1990–1998 (2017) 19. S. Esmaeili, X. Wang, D.J. Lloyd, W.J. Poole, On the precipitation hardening behavior of the Al-Mg-Si-Cu alloy AA6111. Metall. Mater. Trans. A 34, 751–763 (2003)

Effect of Zr Content on the Microstructure, Mechanical Properties, and Corrosion Resistance of Ti–27Nb–xZr Alloys Ying Xu, Huanhuan Wang, Yanqing Cai and Ziyan Wei

Abstract In order to improve comprehensive mechanical properties and corrosion resistance of Ti–Nb–Zr alloys, Ti–27Nb–xZr (0–10 wt%) alloys were prepared by powder metallurgy (PM) method. The effect of Zr content on microstructure, mechanical properties, and corrosion resistance was researched. It was observed that the alloys possessed equal-axis β phase and a little acicular α phase. The mechanical property test showed that with the increase of Zr content, the elastic modulus of Ti–27Nb–xZr alloys was reduced and the compressive strength was improved. When the Zr content was 6 wt%, the Ti–27Nb–6Zr alloy had the highest compressive strength of 625 MPa, and the lowest elastic modulus of 50 GPa. Corrosion resistance of the Ti–27Nb–xZr alloys was evaluated by the potentiodynamic polarization curves. The result showed that the Ti–27Nb–6Zr alloy had better corrosion resistance than those of other alloys with different Zr contents, which exhibited a great potential for orthopedic applications.

1 Introduction Titanium alloys were widely used as implant materials due to their excellent biocompatibility, corrosion resistance, and moderate mechanical properties [1–4]. For example, Ti–6Al–4V alloys were always introduced and used for biomedical applications [5]. However, some problems still exist about Ti–6Al–4V alloys. First, it can release the highly toxic V and Al into the cell tissue during long-term implantation into the human environment, which could cause allergic reactions [6]. Second, Ti–6Al–4V alloys as the implants could easily cause the problem of stress shielding and result in bone resorption due to their higher elastic modulus than that of natural bone [7]. Thus, in order to solve the above problem and achieve better biocompatibility, it has attracted much attention that new titanium alloys are prepared to Y. Xu (B) · H. Wang · Y. Cai · Z. Wei College of Material Science and Engineering, North China University of Science and Technology, Tangshan, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_43

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decrease elastic modulus using to nontoxic elements. Semlitsch et al. researched that the biocompatibility of Ti–6Al–7Nb alloy was not only improved, but also mechanical performance was equivalent with Ti–6Al–4V alloy [8]. However, it was not still solved that Ti–6Al–7Nb and Ti–6Al–4V alloy had high elastic modulus. Niinomi reported that the elastic modulus of alloys has been extensively reduced by the formation of β phase [9]. Some studies reported that mechanical properties and corrosion resistance were improved by adding to β stabilizing elements with nontoxic, including Ta, Nb, Zr, Sn, and Mo [10–12]. Ivasishin et al. found that the β microstructure had better corrosion resistance and lower elastic modulus than α phase, enhancing tissue response [13]. Thus, low-modulus and nontoxic titanium alloys with β phase have attracted more and more attention. Nb is a strong β-stabilizer and no allergic reaction contributing to decreasing the elastic modulus and improving the biocompatibility [14–16]. In addition, Zr is a neutral element with excellent biocompatibility and it is beneficial to increase the compressive strength of the alloy due to solution strengthening [17, 18]. Yu et al. [19] reported that Ti–25Nb–3Zr alloy exhibited the maximum tensile strength about 775 MPa and the lowest elastic modulus about 62 GPa. Guo et al. [20] claimed that Ti–24Nb–4Zr–7.9Sn alloy that was prepared by Powder Metallurgy (PM) method at 1250 °C for 2 h exhibited a lower elastic modulus (about 63.5 GPa), moderate strength, and good corrosion resistance. As mentioned above, the elastic modulus of alloys is still relatively higher than natural bone. Thus, there are three aspects to consider for the preparation of biomedical titanium alloy. First, it is necessary to decrease effectively the elastic modulus comparing with the human bone tissues. Second, the better composition ratio of Nb and Zr in titanium matrix is contribute to the reduction of elastic modulus, the enhance of biocompatibility and corrosion resistant. Finally, in order to resist the external force and extend the service life of material, compression strength should be improved. In the present work, using titanium hydride powder, niobium powder, and zirconium powder as raw material, the Ti–27Nb–xZr alloys containing 0, 2, 4, 6, 8, and 10 wt% Zr were prepared by the powder metallurgy (PM) method. It was investigated that the Zr content of alloys had an influence on the microstructure, phase composition, mechanical properties, and corrosion resistant.

2 Experimental Method 2.1 Preparation of Ti–27Nb–XZr Alloys Titanium hydride powder, niobium powder, and zirconium powder were blended using to stainless steel vacuum ball grinding tank by planet ball mill, and it is used to produce Ti–27Nb–xZr alloys containing 0, 2, 4, 6, 8, and 10 wt% Zr. The alloys were prepared by powder metallurgy method. Green compacts were formed under uniaxial pressure of 400 MPa and were sintered in a high-temperature vacuum

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furnace (10−3 Pa). The sintering process consists of three stages. First, the sintering process started from room temperature up to 200 °C at a heating rate of 5 °C/min drying the green compacts and avoiding cracking problems. Then, the temperature was increased to 800 °C for 2 h, which was used in the dehydrogenation of titanium hydride. Finally, the temperature was raised up to 1400 °C for 2 h, and then the green compacts were furnace-cooled to room temperature.

2.2 Microstructure Characterization The density and porosity of samples were measured by water displacement method using Archimedes principle. The microstructure and phase constitutions were analyzed by Olympus BX51M light optical microscope and X-ray diffractometry (XRD). The sintered samples were grinded with SiC papers of different granulometry, and polished with an alumina-based solution (1 μm) and fine polishing with silica gel. The samples for microstructural analysis were etched using Kroll reactant (5 ml HF + 10 ml HNO3 + 85 ml H2 O).

2.3 Mechanical Properties The samples for measuring the mechanical properties were cylinders (F10 × 6 mm). Elastic modulus (E) and compressive strength were measured at a speed of 1 mm/min by using a WDW-200 machine.

2.4 Electrochemical Tests Corrosion resistance was characterized by the electrochemical tests in a simulated body fluid (SBF). Measurements were carried out in a conventional three-electrode cell with an Ag/AgCl as a reference electrode and a platinum sheet as an counter electrode using an Princeton VersaSTAT 4. After open-circuit potential (OCP) was measured for 30 min, potentiodynamic polarization curves were carried out from − 0.5 to 0.5 V at a scan rate of 0.5 mV/s. The electrochemical parameters, including the corrosion potential (E corr ), the corrosion current density (icorr ) determined by the Tafel slope extrapolation, and the passive current density (ipass ) were obtained from the anodic polarization curves. Simulated body fluid (SBF) was prepared using NaCl, KCl, K2 HPO4 , CaCl2 ·2H2 O, MgCl2 ·6H2 O, NaHCO3 , and Na2 SO4 [21]. The pH was adjusted with 0.1 mol/L HCl and tris (hydroxymethylaminomethane) to 7.25 (36.5 °C).

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3 Results and Discussion 3.1 Microstructure and Phase Composition Figure 1 shows the microstructures of the Ti–27Nb–xZr alloys. A β-matrix and a little needle-like α phase of alloys were produced. It was observed from Fig. 1a–d that with the increase of Zr, the needle-like α phase was transformed into homogeneous equiaxial β phase. When the content of Zr was 6 wt%, the microstructures of alloy had a uniform particle size of 25 μm and clear grain boundaries. It means that Zr also stabilizes beta titanium alloy although Zr was widely considered as a neutral element. It may be caused that the increased Zr content decreased the β phase transus temperature. As shown in Fig. 1d–f, particle size was gradually decreased due to the increase of Zr content. This also indicated the effects of Zr to stabilize β phase and refine the grain size of the alloys. In order to study the phase transformation of alloys, XRD patterns of sintered Ti–27Nb–xZr alloys are exhibited in Fig. 2. It was seen that the diffraction peaks of α + β phase were predominant and the diffraction peaks of Nb and Zr did not appear, owing to full the diffusion and dissolution. It indicated that diffraction peak corresponding to β titanium was significantly enhanced when increasing the content of Zr to 6 wt%, which proved that Ti–27Nb–6Zr alloy had a predominance of β phase. The evolution of porosity and sintered density was measured by water displacement method. The sintered density and porosity of the samples are shown in Fig. 3. It can be seen that the density of the Ti–27Nb–xZr alloys increased with the increase of Zr content, conversely, the total porosity decreased. The alloy with 6 wt% Zr had the highest sintered density of 4.12 g/cm3 and the porosity of 20%. When Zr content was less than 6%, the amount of liquid phase increased with the increase of Zr content by liquid phase sintering at high temperature. The increase of liquid phase sintering was beneficial to the diffusion of atoms, which increased the density of the alloy. When the Zr content was more than 6%, most of Zr still existed in the liquid phase and agglomerated into a sphere under the action of surface tension with the rapid increase of Zr content, and secondary phase particles were formed after sintering. The increase of Zr was conducive to producing the more secondary-phase particles in the alloy. The secondary phase particles inhibited the diffusion of grain boundary, which caused the higher porosity and the lower density of the sintered alloy. Thus, the density of sintered Ti–27Nb–xZr alloys changed as shown in Fig. 3. Figure 1d displays that pore is closed and isolated. It can be seen that slender opening pores evolved into oval closed pores with the increase of Zr contents.

3.2 Mechanical Properties Figure 4 shows stress–strain curves of Ti–27Nb–(0–10)Zr alloys obtained by compression test at room temperature. The compressive strength and elastic modulus of

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Fig. 2 XRD patterns of sintered Ti–27Nb–xZr alloys with different Zr contents

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Ti–27Nb–xZr alloys with different Zr contents sintered at 1400 °C for 2 h are shown in Fig. 5. In Fig. 4, all alloys exhibit single-stage yielding. It can be indicated that the Ti–27Nb–xZr alloys show similar mechanical behavior. Furthermore, alloys first underwent an elastic deformation and then started to deform plastically. According to Fig. 4, the deformation of alloys did not increase monotonously with the increase of Zr content and exist obvious fluctuations due to the homogenous of microstructure. When Zr content was 0%, the compressive deformation was the smaller about 9.3%. The largest compressive deformation of 11.8% was obtained in the Ti–27Nb–6Zr alloy, indicating the better elasticity. The strength first increased with increasing Zr content, until it reached 6 wt% and then it began to decrease. This indicated that the addition of Zr caused solid-solution hardening. When Zr content was 0 wt%, the corresponding compressive strength was the higher up to 417 MPa. When Zr content was increased to 6 wt%, the compressive strength was increased to 625 MPa. When Zr content was increased from 0 to 6 wt%, the elastic modulus was increased from 39 to 50 GPa. When Zr content was increased from 6 to 10 wt%, the compressive strength decreased to 495 MPa and the increase of elastic modulus was not obvious. Zr atom may be soluble in the bcc-structured β phase according to the atomic sizes of Ti, Nb, and Zr [19]. Thus, up to 6 wt% of Zr content is within the solubility limit in the Ti–27Nb–xZr alloys, and Ti–27Nb–6Zr alloy has the lower elastic modulus and higher strength. It avoids stress shielding, which prevents bone resorption in orthopedic implants applications.

3.3 Electrochemical Characterization Apart from having desirable mechanical properties, it is also critical that the implanted biomaterials have excellent corrosion resistance in a body fluid environment [22]. Figure 6 shows the potentiodynamic polarization curves of sintered Ti–27Nb–xZr alloys with different Zr contents. It can be found that the potentiody-

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namic polarization curves of sintered Ti–27Nb–xZr alloys with different Zr contents are similar. It was observed that alloys exhibited similar the process of redox of cathode and the activation–passivation behavior of anode, indicating spontaneous passivation behavior during the test. Thus, the oxide films were spontaneously developed at the surface of the alloys when alloys were immersed in the SBF. The corrosion parameters obtained from the polarization curves are shown in Table 1. The results showed that the corrosion resistance tendency of sintered Ti–27Nb–xZr alloys with different Zr contents was improved. When Zr content was 6 wt%, E corr was reached to the highest value. When Zr content was more than 6 wt%, the E corr began to decrease. The values of E corr , icorr , and ipass indicated that the Ti–27Nb–6Zr alloy had the lowest icorr (about 102.262 μA/cm2 ) and ipass (about 2.50 × 10−4 A/cm2 ), while the highest E corr (about −255.194 mV). The lowest icorr and ipass values indicated that a stable protective film was formed under immersion into the electrolyte. The environment, composition, microstructure, and porosity of titanium alloys have a synthetic influence on the corrosion resistance of titanium alloys. Guo et al. reported

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Fig. 6 Potentiodynamic curves of sintered Ti–27Nb–xZr alloys with different Zr contents

2%

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Potential/V Table 1 Electrochemical parameters of sintered Ti–27Nb–xZr alloys with different Zr contents

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ipass (A/cm2 )

Ti27Nb

−325.249

197.012

6.30 × 10−4

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−262.744

168.384

4.46 × 10−4

Ti27Nb4Zr

−283.874

113.012

4.51 × 10−4

Ti27Nb6Zr

−255.194

102.262

2.50 × 10−4

Ti27Nb8Zr

−341.855

130.786

5.62 × 10−4

Ti27Nb10Zr

−350.963

146.576

5.37 × 10−4

that the more porous materials had a larger surface area exposed to the SBF, which was more vulnerable to corrosion than the less porous materials [20]. Ti–27Nb–6Zr alloy had higher density and lower porosity, and then indicated a higher corrosion resistance. Min et al. [23] claimed that the continuous β matrix could enhance the stability of the passive film, indicating high corrosion resistance. Ti–27Nb–6Zr alloy had homogeneous and continuous equiaxial β-dominant phase with the clear grain boundaries, which indicated high corrosion resistance. The main reason may be that the interface between the grains of continuous equiaxial β phase forms a dense structure in the form of a grain boundary, which may restrain effectively corrosion caused by intergranular voids and defects. Therefore, Ti–27Nb–6Zr alloy can be a potential implant material due to its good corrosion resistance and excellent comprehensive mechanical properties.

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4 Conclusions Ti–27Nb–xZr alloys has been developed by the PM method sintering at 1400 °C for 2 h. Microstructure, mechanical properties and corrosion resistance of alloys were investigated. The conclusions are summarized as follows: (1) The alloys have β-matrix and a little needle-like α phase. Ti–27Nb–6Zr alloy with the uniform particle size of 25 μm has homogeneous equiaxial β-dominant phase and clear grain boundaries, revealing that Zr also stabilizes beta titanium alloy. (2) The addition of Zr improves the compressive strength and reduces the elastic modulus of Ti–27Nb–xZr alloys. Ti–27Nb–6Zr alloy has the highest compressive strength of 625 MPa, however, its elastic modulus is only 50 GPa. (3) Ti–27Nb–6Zr alloy has the lowest icorr value of 102.262 μA/cm2 and the highest E corr value of −255.194 mV, indicating a good corrosion resistance. Acknowledgements This study was financially supported by the Provincial Natural Science Foundation and Key Basic Research Project of Hebei Province (No. C2018209270).

References 1. M. Niinomi, Recent metallic materials for biomedical applications. Metall. Mater. Trans. A 33(3), 477 (2002) 2. Y. Yao, X. Li, Y.Y. Wang, W. Zhao, G. Li, Microstructural evolution and mechanical properties of Ti-Zr β titanium alloy after laser surface remelting. J. Alloys. Compd. 583, 43–47 (2014) 3. F.A. Müller, M.C. Bottino, L. Müller, A.R. Vinicius, In vitro apatite formation on chemically treated (P/M) Ti-13Nb-13Zr. Dent. Mater. 24(1), 50–56 (2008) 4. L. Wang, W. Lu, J. Qin, F. Zhang, D. Zhang, Influence of cold deformation on martensite transformation and mechanical properties of Ti–Nb–Ta–Zr alloy. J. Alloys. Compd. 469(1–2), 512–518 (2009) 5. L. Bolzoni, E.M. Ruiz-Navas, E. Gordo, Evaluation of the mechanical properties of powder metallurgy Ti-6Al-7Nb alloy. J. Mech. Behav. Biomed 67, 110 (2016) 6. Y. Okazaki, Y. Ito, K. Kyo, T.: Corrosion resistance and corrosion fatigue strength of new titanium alloys for medical implants without V and Al. Mat. Sci. Eng. A 213(1–2), 138–147 (1996) 7. R. Huiskes, H. Weinans, R.B. Van, The relationship between stress shielding and bone resorption around total hip stems and the effects of flexible materials. Clin. Orthop. Relat. Res. 274(274), 124–134 (1992) 8. M. Niinomi, Metallic biomaterials. J. Artif. Organs 11(3), 105 (2008) 9. Y.L. Zhou, D.M. Luo, Corrosion behavior of Ti-Mo alloys cold rolled and heat treated. J. Alloys. Compd. 509(21), 6267–6272 (2011) 10. M. Semlitsch, H. Weber, R. Steger, 15 Jahre Erfahrung mit Ti6AI7Nb-Legierung für Gelenkprothesen—Fifteen years of experience with a Ti6AI7Nb alloy for joint replacements. Biomed. Eng. 40(12), 347–355 (1995). Online 11. L.J. Xu, Y.Y. Chen, Z.G. Liu, F.T. Kong, The microstructure and properties of Ti-Mo-Nb alloys for biomedical application. J. Alloys. Compd 453(1–2), 320–324 (2008) 12. Y. Sasikumar, N. Rajendran, Surface modification and in vitro characterization of Cp-Ti and Ti-5Al-2Nb-1Ta alloy in simulated body fluid. J. Mater. Eng. Perform. 21(10), 2177–2187 (2012)

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13. O.M. Ivasishin, P.E. Markovsky, Y.V. Matviychuk, S.L. Semiatin, C.H. Ward, S. Fox, A comparative study of the mechanical properties of high-strength β-titanium alloys. J. Alloys Compd. 457(1), 296–309 (2008) 14. M. Göttlicher, M. Rohnke, A. Helth, T. Leichtweiß, T. Gemming, A. Gebert, J. Eckert, J. Janek, Controlled surface modification of Ti-40Nb implant alloy by electrochemically assisted inductively coupled RF plasma oxidation. Acta Biomater. 9(11), 9201–9210 (2013) 15. A. Terayama, N. Fuyama, Y. Yamashita, I. Ishizaki, H. Kyogoku, Fabrication of Ti-Nb alloys by powder metallurgy process and their shape memory characteristics. J. Alloys Compd. 577(1), S408–S412 (2013) 16. J.M. Chaves, O. Florêncio, P.S. Silva Jr., P.W.B. Marques, S.G. Schneider, Anelastic relaxation associated to phase transformations and interstitial atoms in the Ti-35Nb-7Zr alloy. J. Alloys Compd. 616, 420–425 (2014) 17. W. Simka, A. Krzala, M. Maselbas, G. Dercz, J. Szade, A. Winiarski, J. Michalska, Formation of bioactive coatings on Ti-13Nb-13Zr alloy for hard tissue implants. RSC. Adv. 3(28), 11195–11204 (2013) 18. Y. Okazaki, A New Ti-15Zr-4Nb-4Ta alloy for medical applications. Curr. Opin. Solid State Mater. Sci. 5(1), 45–53 (2001) 19. Y. Zhou, Y. Li, X. Yang, Z. Cui, S. Zhu, Influence of Zr content on phase transformation, microstructure and mechanical properties of Ti75-x Nb25 Zrx (x  0 – 6) alloys. J. Alloys Compd. 486(1), 628–632 (2009) 20. S. Guo, A. Chu, H. Wu, C. Cai, X. Qu, Effect of sintering processing on microstructure, mechanical properties and corrosion resistance of Ti-24Nb-4Zr-7.9Sn alloy for biomedical applications. J. Alloys Compd. 597(6), 211–216 (2014) 21. C.G. Ágreda, M.W.D. Mendes, J.C. Bressiani, A.H.A. Bressiani, Apatite coating on titanium samples obtained by powder metallurgy. Adv. Sci Technol. 86, 28–33 (2013) 22. G.H. Lv, H. Chen, L. Li, E.W. Niu, H. Pang, B. Zou, S.Z. Yang, Investigation of plasma electrolytic oxidation process on AZ91D magnesium alloy. Curr. Appl. Phys. 9(1), 126–130 (2009) 23. X.H. Min, S. Emura, N. Sekido, T. Nishimura, K. Tsuchiya, K. Tsuzaki, Effects of Fe addition on tensile deformation mode and crevice corrosion resistance in Ti-15Mo alloy. Mater. Sci. Eng. A 527(10–11), 2693–2701 (2010)

Preparation of NiO by Precipitation Transformation and Its Supercapacitor Performance Xinglei Wang, Yunqing Liu, Fan Zhang and Xiuling Yan

Abstract Ni(OH)2 precursor was prepared by precipitation and transformation method using nickel nitrate and sodium oxalate as raw materials and cetyltrimethylammonium bromide (CTAB) as dispersant. After heat treatment, the obtained NiO powder was characterized by X-ray diffraction (XRD), scanning electron microscope (SEM), cyclic voltammetry (CV), constant current charge-discharge method, and electrochemical impedance spectroscopy (EIS) techniques. The effects of surfactant concentration and calcination temperature on the morphology, structure, and electrochemical properties of NiO were studied. The results show that the NiO obtained under the condition of CTAB concentration of 0.8 mol/L and heat treatment temperature of 300 °C presents a typical pseudocapacitance property in the potential range of 0–0.4 V (vs. SCE). Under the operating voltage of 0.4 V and the current density of 5 mA/cm2 , the specific capacitance of the single electrode is 536 F/g. By optimizing the preparation conditions, the single-electrode specific capacitance of the NiO material has been significantly improved.

1 Introduction With the booming of electric vehicles, mobile portable consumer devices, and the serious environmental problems caused by traditional energy consumption, the demand for new energy storage, and conversion equipment is becoming more and more urgent. As a new electrical energy storage and conversion device, supercapacitors have many advantages and their potential application value and huge market prospect have attracted wide attention of researchers all over the world [1–8]. The electrode material is the core of the supercapacitor and its structure and properties have great influence on the performance of supercapacitors. Therefore, it is particularly important to study the supercapacitor electrode materials to obtain electrode materials with better performance, to realize the widespread application of superX. Wang · Y. Liu · F. Zhang · X. Yan (B) Laboratory of Condensed Matter Phase Transition and Microstructure, School of Chemistry and Environmental Sciences, Yili Normal University, Xinjiang, Yining, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_44

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capacitors, and to change the current energy consumption pattern. As a tantalum capacitor electrode material, NiO is considered to be a potential application electrode material [9–16] because of its high theoretical specific capacitance, low-cost, abundant resources, and easy preparation. Therefore, research on supercapacitor electrode materials has obtained an electrode material with superior performance and the focus is on studying the capacitance and cycle performance of the nickel oxide electrode material to realize the demand for preparing a low-cost and highcapacity supercapacitor. The theoretical capacity of nickel oxide is relatively large (up to 2584 F/g in the electrochemical window of 0.5 V), which is of great value for development. There have been several reports on nickel oxide. Among them, Zhang et al. [12] synthesized NiO nanoparticles by simple liquid phase method and calcined at 300 °C to form NiO electrodes with a single electrode specific capacity of 300 F/g. Wang [17] uses nickel nitrate as raw material, SBA-15 as template, and 550 °C heat treatment to obtain ordered porous structure of NiO with a specific capacitance of up to 120 F/g. Liu et al. [18] prepared precursors by using nickel nitrate as raw material, urea as precipitant, and polyethylene glycol (PEG) as a template for hydrothermal synthesis. After heat treatment, the shape of the hedgehog nickel oxide was obtained. The specific capacitance of NiO obtained by heat treatment of 300 °C is 290 F/g. In this paper, NiO powder was prepared by precipitation conversion method, and the preparation conditions were further optimized to achieve the purpose of improving the electrochemical capacitance performance of the material.

2 Experimental Part 2.1 Raw Materials and Reagents Nickel nitrate (Ni(NO3 )2 ·6H2 O), cetyltrimethylammonium bromide (C19 H42 BrN, CTAB), sodium oxalate (Na2 C2 O4 ), sodium hydroxide (NaOH), all of which are analytically pure (AR) reagents.

2.2 Preparation of NiO After adding 225 mL of 0.1 mol/L NaC2 O4 solution into 200 mL of 0.1 mol/L Ni(NO3 )2 ·6H2 O solution, 50 mL of a certain concentration of CTAB was added. After mixing, the mixture was stirred for 10 min, then 325 mL of 0.2 mol/L NaOH solution was added and stirred for 2 h. After repeated centrifugation and washing, the obtained product was dried at 80 °C for 12 h, and then NiO powder samples were obtained by calcining at different temperatures for 3 h and cooling. The reaction to producing Ni(OH)2 during the preparation process is as follows:

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Ni(NO3 )2 + Na2 C2 O4 → NiC2 O4 + 2NaNO3 NiC2 O4 + 2NaOH → Ni(OH)2 + Na2 C2 O4 The reaction is carried out by precipitation conversion method, and Ni(NO3 )2 and Na2 C2 O4 are reacted to form precipitated NiC2 O4 , which is further transformed under NaOH basic conditions. Precipitated with Ni(OH)2 to obtain the precursor.

2.3 Physical Characterization and Performance Testing of Samples XRD measurement was carried out using a German Duke D8 ADVANCE A25 X-ray diffractometer. The experimental conditions were: Cu Kα radiation (λ  1.54060 Å), tube voltage 20 kV, tube current 5 mA, scanning range 2θ  20° − 80°, scanning speed was 10°/min. The morphology and particle size of the sample were observed with a JSM-7500F scanning electron microscope from Japan.

2.4 Preparation and Electrochemical Testing of NiO Electrode The prepared NiO powder is mixed with acetylene black, polytetrafluoroethylene (PTFE) and carboxymethyl cellulose (CMC) to form a paste, uniformly coated on a nickel mesh as a working electrode, and dried at normal temperature. It was post pressed into an electrode sheet having an area of 1 cm2 . The electrochemical test was performed by the CHI 660D (Shanghai) electrochemical workstation. The three electrode system, the saturated calomel electrode (SCE) was used as the reference electrode, and the platinum electrode was used as the auxiliary electrode. The mass of the electroactive substance was 5 mg. Cyclic voltammetry and constant current charge and discharge tests were carried out in a 5 mol/L KOH solution and a potential range of 0–0.4 V.

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3 Results and Discussions 3.1 Instrument Characterization and Analysis of Electrode Materials 3.1.1

XRD Analysis of Electrode Materials

Figure 1 is an XRD pattern of NiO prepared at different calcination temperatures. It can be seen from the figure that the four patterns are basically the same, and there are obvious diffraction peaks. The characteristic diffraction peaks of 2θ are 37.2°, 43.2°, and 62.9° respectively correspond to the (111), (200), and (220) diffraction planes of the cubic system. And consistent with the standard map JCPDS (73-1523), no other impurity peaks exist, fully indicating that the product obtained by calcination is cubic crystal NiO. The diffraction peak is wider when the temperature is 250 °C, but the peak intensity is enhanced when the temperature is up to 350 °C, which indicates that the product has been transformed to crystal form at this time. The comparison shows that the positions of the peaks are basically the same at different calcination temperatures, and the change of calcination temperature does not change the crystal form of the material. However, the intensity of the diffraction peak changes with the change of calcination temperature: When the calcination temperature is 250 °C, the diffraction peak intensity is the weakest, not sharp enough, and there are many heterogeneity peaks, which indicates that Ni(OH)2 is not completely oxidized; When the temperature is 400 °C, the diffraction peak is the sharpest, the half width is the narrowest, the crystallinity is the best, and the impurity peak is the least; when the calcination temperature is between 300 and 350 °C, the intensity of the diffraction peak and the sharpness are in the middle. As the calcination temperature increases, the material gradually crystallizes, the crystal grains grow gradually, and the surface area of the material decreases. According to the peak intensity, it can be judged the product obtained by 300 °C is amorphous, and the amorphous material is suitable for use as a supercapacitor electrode. The material has good electrochemical capacitance performance [17]. Therefore, the NiO powder obtained at the calcination temperature of 300 °C is selected as the research object. The electrochemical properties of NiO prepared at different calcination temperatures are tested and the chemical capacitance performances are studied to determine the optimum calcination temperature.

3.1.2

Effect of Surfactant Concentration on the Morphology of NiO

In Fig. 2a–d are SEM images of NiO materials prepared by surfactant concentrations of 0.1, 0.5, 0.8, and 1.0 mol/L, respectively. For the capacitor electrode material, the loose structure represents a large surface area, and thus has a large active surface area during the electrode reaction, thereby having a strong reactivity. It can be seen from the figure that the crystal is agglomerated and the structure is incomplete; in Fig. 2b, the structure is loose and porous, but there is still agglomeration; in Fig. 2c,

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Fig. 1 XRD patterns of NiO prepared at 250 °C (a), 300 °C (b), 350 °C (c) and 400 °C (d)

the development of the grain is gradually completed, and the sheet structure is clear; the obvious agglomeration phenomenon can be seen in Fig. 2d. The possible reason is that the crystal is bound more and more closely, which reduces the surface area and affects the electrochemical properties of the materials. With the increasing concentration of surfactants, the development of crystal grains is gradually completed, and the particle size is getting larger and larger. When the surfactant concentration is 0.8 mol/L, the prepared SEM of NiO shows a clear NiO sheet structure, which has a large surface area due to porous and provides an important morphological basis for the specific capacity of the electrode material. Ni(OH)2 precursor was prepared by precipitation conversion method. The function of CTAB was to coat the surface of Ni(OH)2 particles which have been formed and to suppress the formation of Ni(OH)2 by Ni ions. The growth rate of the surface of the particles controls the size of the Ni(OH)2 particles; on the other hand, the generated Ni(OH)2 particles are isolated from each other, inhibiting the occurrence of agglomeration, i.e., the concentration of the surfactant directly affects dispersibility and particle size to Ni(OH)2 . In order to further understand the NiO material prepared under the same calcination temperature (300 °C) and different concentration (CTAB) conditions, the following electrochemical performance tests were carried out.

3.2 Electrochemical Performance Test Results and Discussion 3.2.1

Cyclic Voltammetry Curve of NiO Electrode

The electrochemical properties of NiO electrode materials prepared by calcination at different concentrations of CTAB and 300 °C were determined by cyclic voltammetry using 5 mol/L KOH as electrolyte. The results are shown in Fig. 3. It can be clearly

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Fig. 2 SEM images of NiO prepared with different concentrations of CTAB: 0.1 mol/L (a), 0.5 mol/L (b), 0.8 mol/L (c) and 1.0 mol/L (d)

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Fig. 3 CV curves of NiO electrode in 5 mol/L KOH at 10 mV/s. The concentrations of surfactant CTAB for the preparation of NiO are: 0.1 mol/L (a), 0.5 mol/L (b), 0.8 mol/L (c) and 1.0 mol/L (d)

seen from the figure that the NiO electrode obtained at four concentrations has a strong redox peak in the 0–0.4 V potential range, showing a significant Faraday quasi-capacitance characteristic. As the concentration of the surfactant increases, the position of the redox peak of the electrode gradually shifts. When the CTAB concentration is 0.8 mol/L, the curve encloses the largest area, the peak current is the strongest, and the capacitance value is the largest. When the CTAB concentration continues to increase to 1.0 mol/L, the redox peak becomes smaller. The possible reason is that the generated particles become larger due to the increased concentration of CTAB, which lowers the specific surface area, resulting in a decrease in peak current. This also confirms the results of Fig. 2. Therefore, 0.8 mol/L is selected as the best surfactant concentration. SEM characterization preliminary shows that 300 °C is the best calcination temperature, and the corresponding electrochemical performance characterization experiments were carried out to verify this. Figure 4 shows the cyclic voltammetry curves of the NiO electrode at a sweep speed of 10 mV/s, which prepared at the CTAB concentration of 0.8 mol/L and the calcination temperature of 250, 300, 350 and 400 °C, respectively. It can be clearly seen from Fig. 4 that the NiO electrode obtained at four temperatures has a strong redox peak in the cyclic voltammetry range of 0–0.4 V. The redox peak of the electrode increases with the calcination temperature. With the increase of calcination temperature, the position of the redox peak changes at 300 °C, indicating that the maximum capacitance is obtained when the calcination temperature is 300 °C, and the peak current at other temperatures is relatively small. The results show that the peak current of the electrode increases at first and then decreases with the increase of calcination temperature. This phenomenon is mainly related to their microstructure: with the increasing of calcination temperature, the agglomeration of NiO crystals occurs, the crystals pile up together, and the surface contact between OH− and NiO active material in the electrolyte decreases, which affects the electrochemical per-

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Fig. 4 The CV curves of NiO electrode in 5 mol/L KOH at scan speed of 10 mV/s. The calcination temperatures for the preparation of NiO are: 250 °C (a), 300 °C (b), 350 °C (c) and 400 °C (d)

Fig. 5 Cyclic voltammograms of NiO electrode in 5 mol/L KOH at different sweep speeds

formance of NiO. The reason for the low current at 250 °C is that Ni(OH)2 is not completely oxidized. When the concentration of CTAB is 0.8 mol/L and the calcination temperature is 300 °C, the prepared NiO is used as the working electrode, and the cyclic voltammetry is performed at different sweep speeds. From Fig. 5, it can be seen that the cyclic voltammograms curve of NiO electrode shows a pair of redox peaks in the range of 0–0.4 V, which indicates that the electrode has typical Faraday quasi-capacitance characteristics. Moreover, with the increase of scanning speed, the pattern does not change greatly, which shows that the stability of the material is better under the condition of high current.

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Constant Current Charge and Discharge Performance Test

Constant current discharge is an important means to test the behavior of the capacitor. The discharge curve can be used to calculate the capacity of the capacitor. Figure 6 shows the galvanostatic charge-discharge test curves of NiO electrodes obtained at different concentrations of CTAB. Its single-electrode mass ratio capacitance C m can be expressed by the following formula: Cm 

Q I × t  U × m U × m

(1)

where Q is the amount of electricity, U is the discharge voltage range, t is the discharge time, I is the charge and discharge current, and m is the amount of electrode material. According to this formula, the mass ratio of the sample electrode can be calculated as shown in Table 1. Comparing the specific capacities of the electrodes at four concentrations, it is known that the NiO electrode material with a CTAB concentration of 0.8 mol/L has the highest specific electrode capacity. As the concentration of the surfactant increases, the specific capacitance of NiO increases first and then decreases. The possible reason is that when the concentration of CTAB is low, the coating effect on the produced Ni(OH)2 particles is poor, and the particles continue to grow to

Fig. 6 Constant current charge–discharge curves of NiO electrodes in 5 mol/L KOH at current density of 5 mA/cm2 . The concentrations of surfactant CTAB for the preparation of NiO are: 0.1 mol/L (a), 0.5 mol/L (b), 0.8 mol/L (c) and 1.0 mol/L (d)

Table 1 Influence of CTAB concentration on the specific capacitance of NiO electrode

Reactive concentration of CTAB (mol/L)

0.1

0.5

0.8

1.0

Specific capacitance (F/g)

130

501

536

460

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some extent, so the particle size is larger, and the NiO particle obtained after the heat treatment is not uniform and the specific capacitance is smaller. As the concentration of CTAB increases, the coating effect on the Ni(OH)2 particles is gradually enhanced, and the growth of Ni2+ on the surface of the particles is hindered, and the final size of the particles is reduced. However, when the concentration of CTAB is too high, the growth rate of Ni2+ on the crystal nucleus was slowed down, and the time required for complete precipitation of Ni2+ increases, leading to the incomplete conversion of nickel oxalate. Figure 7 shows the constant current charge-discharge curves of a NiO electrode at a current density of 5 mA/cm2 , which prepared at different calcination temperatures with a CTAB concentration of 0.8 mol/L. It can be seen from the shapes of the four curves in Fig. 7 that the charging and discharging platform appears on the charge–discharge curve, which corresponds to the redox peak of the electrode active material, indicating that the specific capacity of the NiO material mainly exhibits the Faraday quasi-capacitance characteristic. The specific capacity of the single electrode calculated by the formula (1) according to the data in Fig. 7 is shown in Table 2. By comparison, the specific capacitance of the electrode obtained by NiO electrode material at 300 °C is the highest. With the further increase of temperature, the agglomeration phenomenon of NiO crystal occurs, which reduces the specific surface area, and leads to the decrease of specific capacity of the electrode. Figure 8 shows the constant current charge-discharge curves of the prepared NiO electrode at different current densities when the CTAB concentration is 0.8 mol/L and the calcination temperature is 300 °C. It can be seen from Fig. 8 that the charge-

Fig. 7 Constant current charge–discharge curves of NiO electrodes in 5 mol/L KOH at current density of 5 mA/cm2 . The calcination temperatures for the preparation of NiO are: 250 °C (a), 300 °C (b), 350 °C (c) and 400 °C (d)

Table 2 Influence of calcination temperature on the specific capacitance of NiO electrode

Calcination temperature (°C)

250

300

350

400

Specific capacitance (F/g)

446

536

466

425

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Fig. 8 Constant current charge–discharge curves of NiO electrodes in 5 mol/L KOH at different current density of 5 mA/cm2 (a), 10 mA/cm2 (b), 15 mA/cm2 (c) and 20 mA/cm2 (d)

Table 3 Influence of current density on capacitance of NiO electrode

Current density (mA/cm2 )

5

10

15

20

Specific capacitance (F/g)

536

450

392

286

discharge curve is not completely symmetrical. This is because when charging and discharging, the electrode adsorbs a large amount of electrolyte ions, causing a certain degree of polarization of the electrolyte ions in the liquid phase diffusion, leading to a slight deviation of the charge-discharge curve, but it is basically linear, with the increases of current density, the charge-discharge time of the electrode is reduced accordingly. The specific capacity of the single electrode is calculated as shown in Table 3. Table 3 indicates that as the charge-discharge current increases, the specific capacitance decreases, which may be caused by the lack of effective utilization of electrode material at a large current.

3.2.3

AC Impedance Performance

Figure 9 shows the AC impedance spectra of a single electrode at a calcination temperature of 300 °C and a voltage of 2 mV. The EIS curve consists of two parts, that is, the semicircular arc of the high-frequency region and the straight part of the low-frequency region. The intersection value of the real axis with curve represents the real impedance of the electrode. It can be seen from the figure that the internal resistance of the electrode material at 0.1, 0.5, 0.8, and 1.0 mol/L is about 0.63, 0.67, 0.75, and 0.7 , respectively. The semicircular arc represents the induced charge transfer impedance associated with the performance of the porous electrode material. The semicircular arcs of the four curves in Fig. 9 are smaller, and the arc of the electrode material obtained at 0.8 mol/L is the smallest. This indicates that the electrode material obtained at this concentration has a quick charge transfer

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Fig. 9 EIS of NiO electrode at a voltage of 2 mV in 5 mol/L KOH. The concentrations of surfactant CTAB for the preparation of NiO are: 0.1 mol/L (a), 0.5 mol/L (b), 0.8 mol/L (c) and 1.0 mol/L (d)

process in the electrolyte, namely, smaller charge transfer impedance. The curve in the low-frequency region represents the diffusion process of the electrolyte and the diffusion of OH− ions into the interior of the NiO electrode. The impedance here is the diffusion impedance. It can be clearly seen from Fig. 9 that the internal resistance of the electrode material obtained at 0.8 mol/L is about 0.75 , and its low-frequency curve almost rises straight, indicating that the ion diffusion speed is faster and the diffusion time is shorter, showing better performance of supercapacitor. Figure 10 is an AC impedance diagram of a prepared NiO single electrode at a voltage of 2 mV with a CTAB concentration of 0.8 mol/L and a different calcination temperature. As can be seen from the figure, the internal resistance of the electrode material is about 0.61, 0.75, 0.65 and 0.7  at 250, 300, 350, and 400 °C, respectively. The semicircular arc in the high-frequency region represents the induced charge transfer impedance associated with the performance of the porous electrode material, and the semicircle arc of the four curves in the figure are smaller, of that 300 °C is the smallest, which indicates that the electrode material obtained at this temperature has a rapid charge transfer process in the electrolyte, that is, a smaller charge transfer impedance. The low-frequency curve of the electrode material obtained under 300 °C almost rises straight, indicating that ion diffusion can be quick, the diffusion time is shortened, and the supercapacitor has better performance. In the high-frequency region, the semicircular radius of the AC impedance curve of the sample is smaller, indicating that the electrode resistance of the electrode electrochemical reaction is smaller. And the intersection of the AC impedance curve and the horizontal axis is small, that is, the liquid junction resistance is small, indicating the conductivity of the electrode material is good. In the low-frequency region, the electrode impedance curve is approximately 45° angle with the coordinate axis, which shows that the proton diffusion performance of the electrode is better, and the angle between AC impedance curve and coordinate axis is approximately 45° at a voltage of 2 mV. The proton diffusion performance of the NiO sample is relatively better at this voltage.

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Fig. 10 EIS of NiO electrode at a voltage of 2 mV in 5 mol/L KOH. The calcination temperatures for the preparation of NiO are: 250 °C (a), 300 °C (b), 350 °C (c) and 400 °C (d)

4 Conclusion In this paper, Ni(OH)2 precursors were prepared by precipitation conversion method with sodium oxalate and nickel nitrate as raw materials, CTAB as surfactant and NaOH as precipitant. The NiO powder was prepared by heat treatment at different temperatures. The effects of the surfactant concentration and the calcination temperature on the electrochemical performance were investigated. The X-ray diffraction and SEM image analyses showed that the prepared NiO agglomerated into loose particles. The results of cyclic voltammetry, constant current charge-discharge and AC impedance presented that the prepared NiO had the Faraday quasi-capacitance characteristics, and the surfactant concentration and the calcination temperature had great influences on the specific capacity. When the concentration of CTAB is 0.8 mol/L, and the calcination temperature is 300 °C, the prepared NiO supercapacitor shows the best performance, and the specific capacity of the electrode material can reach 536 F/g under the charge–discharge current density of 5 mA/cm2 . By optimizing the preparation conditions, the single electrode specific capacitance of NiO material is significantly improved. Acknowledgements This work is supported by the National Natural Science Foundation of China (51561030).

References 1. B.E. Conway, Ransition from “supercapacitor” to “battery” behavior in electrochemical energy storage. J. Electrochem. Soc. 138(6), 1539–1548 (1991) 2. G.Y. Zhao, C.L. Xu, H.L. Li, Highly ordered cobalt-manganese oxide (CMO) nanowire array thin film on Ti/Si substrate as an electrode for electrochemical capacitor. J. Power Sources

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163(2), 1132–1136 (2007) 3. R.A. Huggins, Supercapacitors and electrochemical pulse sources. Solid State Ionics 134(1–2), 179–195 (2000) 4. D.Y. Zhang, M. Han, Y.B. Li et al., Fabrication of the nitrogen doped ordered porous carbon derived from amino-maltose with excellent capacitance performance. J. Porous Mat. 25, 29–35 (2018) 5. Y.B. Yang, P.X. Li, S.T. Wu et al., Hierarchically designed three dimensional macro/mesoporous carbon frameworks for advanced electrochemical capacitance storage. Chem. Eur. J. 21, 6157–6164 (2015) 6. J.P. Zheng, T.R. Jow, A new charge storage mechanism for electrochemical capacitors. J. Electrochem. Soc. 142(1), L6–L8 (1995) 7. X.L. Li, W.J. Li, X.Y. Chen et al., Hydrothermal synthesis and characterization of orchid-like MnO2 nanostructures. J. Cryst. Growth 297(2), 387–389 (2006) 8. G. Arabale, D. Wagh, M. Kulkarni et al., Enhanced supercapacitance of multiwalled carbon nanotubes functionalized with ruthenium oxide. Chem. Phys. Lett. 376(1–2), 207–213 (2003) 9. X.M. Liu, Y.H. Zhang, X.G. Zhang et al., Studies on Me/Al-layered double hydroxides (Me  Ni and Co) as electrode materials for electrochemical capacitors. Electrochim. Acta 49(19), 3137–3141 (2004) 10. K.X. He, X.G. Zhang, J. Li, Preparation and electrochemical capacitance of Me double hydroxides (Me  Ni and Co)/TiO2 nanotube composites electrode. Electrochim. Acta 51, 1289–1292 (2006) 11. B. Parama, E.L. Derrek, M. Rick, Q. Zhang et al., Electrochemical capacitance of Ni-doped metal organic framework and reduced graphene oxide composites: more than the sum of its parts. ACS Appl. Mater. Inter. 7(6), 3655–3664 (2015) 12. F.B. Zhang, Y.K. Zhou, H.L. Li, Nanocrystalline NiO as an electrode material for electrochemical capacitor. Mater. Chem. Phys. 83, 260–264 (2004) 13. K.W. Nam, W.S. Yoon, K.B. Kim, X-ray absorption spectroscopy studies of nickel oxide thin film electrodes for supercapacitors. Electrochim. Acta 47, 3201–3209 (2002) 14. V. Ganesh, V. Lakshminarayanan, S. Pitchumani, Assessment of liquid crystal template deposited porous nickel as a supercapacitor electrode material. J. Solid State Electrochem. 8(6), A308–A312 (2005) 15. W. Xing, F. Li, Z.F. Yan, G.Q. Lu, Synthesis and electrochemical properties of mesoporous nickel oxide. J. Power Sources 134, 324–330 (2004) 16. Q.H. Huang, X.Y. Wang, J. Li et al., Nickel hydroxide/activated carbon composite electrodes for electrochemical capacitors. J. Power Sources 164(1), 425–429 (2007) 17. D.B. Wang, C.X. Song, Z.S. Hu et al., Fabrication of Hollow Spheres and Thin Films of Nickel Hydroxide and Nickel Oxide with Hierarchical Structures [J]. J. Phys. Chem. B. 109, 1125–1129 (2005) 18. X.M. Liu, X.G. Zhang, S.Y. Fu, Preparation of urchinlike NiO nanostructures and their electrochemical capacitive behaviors. Mater. Res. Bull. 41(3), 620–627 (2006)

Study on Fabrication and Compressive Properties of Mg/Al-Ordered Structure Composites Han Wang, Yu Fu, Mingming Su and Hai Hao

Abstract As a special kind of gas/metal matrix composites with potential application prospects, porous metals have high specific strength integrated with special functional properties. However, since the gas phase in porous metals hardly contributes to the absolute mechanical properties, the properties of porous metals still need to be improved. In this study, Mg/Al-ordered structure composites were prepared successfully by infiltrating the commercial pure magnesium into the ordered porous aluminum to replace the gas phase with the purpose of improving compressive properties. Aiming at decreasing casting defects, the infiltration process of Mg/Al-ordered structure composites was simulated by the software ProCAST to optimize the infiltration temperature and preheated temperature. The results of the quasi-static compression test indicated that the stress-strain curves of Mg/Al-ordered structure composites were similar to the porous metals. There was still stress plateau on the stress-strain curve. Furthermore, the compressive strength and plateau stress were improved obviously compared with the ordered porous aluminum. During the compression deformation, the ordered porous aluminum acted as the skeleton with great plasticity and ductility while the commercial pure magnesium as filler expressed lightweight and high strength. Co-continuous commercial pure aluminum and magnesium with ordered structure led to excellent compressive properties of Mg/Al-ordered structure composites.

1 Introduction As a special kind of gas/metal composite materials, porous metals express lightweight, high specific strength, and some special functional properties. These superior properties help porous metals play an increasingly important role in many industrial fields [1]. However, the absolute mechanical properties of porous metals are lower compared with traditional metallic materials and still need to be improved. TraH. Wang · Y. Fu · M. Su · H. Hao (B) School of Materials Science and Engineering, Dalian University of Technology, Dalian 116023, China e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_45

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ditional methods to improve the mechanical properties of porous metals are focused on changing porous structure including cell shape [2], size [3], and distribution [4]. Strengthening matrix materials by alloying [5, 6] and heat treatment [7, 8] is another effective method. Cheng and Han [9] proposed a view that the gas as the second phase in the porous metals hardly affected the overall mechanical properties. If the gas phase is substituted by a solid phase, the effect of the second phase can be significant and should not be neglected. Therefore, filling the gas phase with a solid phase is a valid method to improve the mechanical properties and composites based on porous metals can be obtained through this method. Kwon et al. [10] filled open-cell foams with elastic rubber. The filled foams had higher compressive strength than foams without filler. Cheng and Han [9] used the silicate rubber to fill open-cell aluminum foams. The plateau region was prolonged and densification strain increased. Yuan et al. [11] infiltrated epoxy resin (ER) into the aluminum foams. The plateau stress and energy absorption capability of the composite structures increased with increasing amount of epoxy resin. Metals as the filler have high strength and stiffness. Li et al. [12] proposed that the intermetallic compound Mg17 Al12 was utilized as a filler. They filled metal foams with the Mg17 Al12 to prepare Mg17 Al12 /Al composite. The compressive properties and energy absorption efficiency were improved. Li et al. [13] investigated the porous titanium with an entangled structure filled with biodegradable magnesium (p-Ti/Mg). It was found that the p-Ti/Mg composite had higher strength than pure magnesium and porous titanium with entangled structure (p-Ti). The literature above are mainly focused on random porous structures. However, ordered porous structure has been a research hotspot recently due to the freedom of structure design and the control of mechanical properties. Esfahani et al. [14] studied a new method for independently tuning the stiffness and toughness of the material by adding various polymers to the additively manufactured porous Ti structures. In this study, the ordered porous aluminum with cubic pores was filled with commercial pure magnesium to prepare Mg/Al-ordered structure composites. The infiltration preparation process of commercial pure magnesium was investigated by software ProCAST and the optimal process parameters including infiltrating temperature and preheated temperature were obtained. The quasi-static compression test was performed to characterize the compressive properties.

2 Experimental Procedures As the matrix of Mg/Al-ordered structure composite, the ordered porous aluminum was prepared by indirect additive manufacturing technique combining the selective laser sintering and the infiltration casting. The matrix material and porosity of ordered porous aluminum were commercial pure aluminum (99.7 wt%) and 54%, respectively. The software ProCAST with finite element method was utilized to simulate the process of filling the ordered porous aluminum with commercial pure magnesium (99.7 wt%). Aiming at decreasing the casting defects, the infiltration

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temperature and preheated temperature were optimized. Mg/Al-ordered structure composites were prepared by infiltration utilizing the optimal process parameters from the simulation. The commercial pure magnesium was melted and held in an electrical resistance furnace with the protection of RJ-2 flux. Then the commercial pure magnesium melt was poured into the preheated mold and infiltrated ordered porous aluminum. After the solidification of commercial pure magnesium melt, the specimens of Mg/Al ordered structure composites were cutting by wire electrode cutting machine. The dimension of the quasi-static compressive specimen was 56 mm × 56 mm × 60 mm. The interface was observed and the compressive properties were characterized by quasi-static compression test at an initial strain speed 10−3 /s.

3 Results and Discussion 3.1 Optimization of Infiltration Parameters In order to decrease the casting defects including insufficient infiltration and shrinkage porosity, the infiltration temperature and preheated temperature were optimized by simulation. In this simulation, pure magnesium was selected as infiltration material and pure aluminum was the matrix of ordered porous aluminum. The mold material was set as stainless steel. The thermal conductivity, density, and other physical parameters of these materials were all obtained from the database of software ProCAST. The effect of infiltration temperature (800, 770, 750, 740, 730 °C) was investigated first and in this situation, the preheated temperature was set at 200 °C. The results of insufficient infiltration and shrinkage are shown in Fig. 1. The blue color represents the area of insufficient infiltration and the purple color represents the area of shrinkage. At higher infiltration temperature (800, 770 °C), the infiltration was sufficient and there was little shrinkage. The commercial pure magnesium is a kind of combustible fillers that cause oxide inclusions adverse to mechanical properties. Therefore, the infiltration temperature above 750 °C was not an optimal temperature for infiltration. At lower infiltration temperature (740, 730 °C), there was obvious insufficient infiltration together with a lot of shrinkages. For lower temperature (750, 740, 730 °C), the preheated temperature was optimized to improve the infiltration process and the results are shown in Fig. 2. The infiltration was sufficient at a higher preheated temperature (250 °C). With preheated temperature decreasing to 200 °C, the phenomena of insufficient infiltration appeared. When the preheated temperature was 150 °C, insufficient infiltration became more obvious regardless of infiltration temperature. The results of shrinkage at different preheated temperatures are shown in Fig. 3. It is found that there was more shrinkage at lower infiltration temperature and preheated temperature. Considering all the simulation results above, the infiltration was easy to be sufficient with little shrinkage when the infiltration temperature and preheated

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Fig. 1 Insufficient infiltration and shrinkage at different infiltration temperatures

Fig. 2 Infiltration results at different infiltration and preheated temperatures

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Fig. 3 Shrinkage at different infiltration and preheated temperatures

temperature were both higher. Therefore, the optimal infiltration process parameters were obtained. The infiltration temperature was 750 °C and the preheated temperature was 250 °C.

3.2 Compressive Properties The quasi-static compression process of ordered porous aluminum is shown in Fig. 4. The main deformation mechanism of ordered porous aluminum was stretching dominated deformation. The struts along with the compression direction were initially compressed. The gas phase hardly resisted the compressive load. There were no obvious macro cracks on the specimen. Therefore, the specimen expressed excellent plasticity and ductility. The quasi-static compression process of Mg/Al-ordered structure composite is shown in Fig. 5. Obvious macro cracks occurred on the middle part of the specimen

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Fig. 4 The quasi-static compression process of ordered porous aluminum in a strain of a 0%, b 10%, c 20%, d 30%, e 40%, and f 50%

Fig. 5 The quasi-static compression process of Mg/Al-ordered structure composite in a strain of a 0%, b 10%, c 20%, d 30%, e 40%, and f 50%

during the compression process. Pure aluminum (FCC crystal structure) had excellent plasticity, while pure magnesium (HCP crystal structure) had inferior plasticity. The interface bonding between magnesium and aluminum was weaker. Therefore, the voids or the interfaces between magnesium and aluminum were the crack sources. With the strain increasing, the cracks propagated quickly and main cracks formed through the whole specimen with a certain angle. The specimen of Mg/Al-ordered structure composite finally fractured with brittle features due to the weak interface and brittle pure magnesium filler. The compressive stress–strain curves of Mg/Al-ordered structure composite and ordered porous aluminum are plotted in Fig. 6. The stress–strain curve of

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Fig. 6 The compressive stress–strain curves of Mg/Al-ordered structure composite and ordered porous aluminum

Mg/Al-ordered structure composite was similar to ordered porous aluminum. The densification strain was determined by the intersection of plateau region and densification region slope. With the strain increasing, the linear elastic stage appeared first. After the specimen yielding, there was still stress plateau on the stress–strain curve. Furthermore, the compressive strength and plateau stress were improved obviously compared with ordered porous aluminum. However, the width of the plateau region and the densification strain decreased compared with ordered porous aluminum. During the compression deformation, the ordered porous aluminum acted as skeleton due to great plasticity and ductility. The commercial pure magnesium as filler showed lightweight and high strength with brittleness features. Co-continuous commercial pure aluminum and magnesium with ordered structure led to excellent compressive properties of Mg/Al-ordered structure composites. The compressive strength, plateau stress, and their specific values (compressive strength to density ratio and plateau stress to density ratio) of Mg/Al-ordered structure composite and ordered porous aluminum are listed in Table 1. It is found that the compressive strength σ c and plateau stress σ p were improved obviously. Considering the effect of density ρ, the compressive strength to density ratio (σ c /ρ) and plateau stress to density ratio (σ p /ρ) were still higher than the specimen without fillers. But the densification strain was a little lower than before. The increase of compressive strength and plateau stress of Mg/Al-ordered structure composite was both beneficial to the improvement of energy absorption capacity. Energy absorption capacities at densification strain were 46.63 MJ/m3 for

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Table 1 The compressive strength σ c , plateau stress σ p , specific compressive strength σ c /ρ, and specific plateau stress σ p /ρ of Mg/Al-ordered structure composite and ordered porous aluminum Specimen

σ c (MPa)

σ c /ρ (MPa g−1 cm3 )

σ p (MPa)

σ p /ρ (MPa g−1 cm3 )

Mg/Al-ordered structure composite

65.06

32.26

122.85

60.91

Ordered porous aluminum

7.01

6.81

25.54

24.82

Mg/Al-ordered structure composite and 12.02 MJ/m3 for ordered porous aluminum. The energy absorption capacity of Mg/Al-ordered structure composite had an obvious advantage over ordered porous aluminum. Considering the energy absorption capacity to density ratio, Mg/Al-ordered structure composites as energy absorbers were still better than ordered porous aluminum.

4 Conclusions (1) The Mg/Al-ordered structure composites were prepared by filling the ordered porous aluminum with commercial pure magnesium. The optimal process parameters including the infiltration casting temperature of 750 °C and the preheated temperature of 250 °C were obtained by simulation. (2) The compressive strength, plateau stress, and energy absorption capacity of Mg/Al-ordered structure composite were improved compared with the ordered porous aluminum. Acknowledgements This work was supported by the National Key Research and Development Program of China (grant numbers 2016YFB0701204).

References 1. J. Banhart, Manufacture, characterisation and application of cellular metals and metal foams. Prog. Mater. Sci. 46(6), 559–632 (2001) 2. B. Jiang, N.Q. Zhao, C.S. Shi, J.J. Li, Processing of open cell aluminum foams with tailored porous morphology. Scr. Mater. 53(6), 781–785 (2005) 3. P. Kenesei, C. Kadar, Z. Rajkovits, J. Lendvai, The influence of cell-size distribution on the plastic deformation in metal foams. Scr. Mater. 50(2), 295–300 (2004) 4. S. He, Y. Zhang, G. Dai, J. Jiang, Preparation of density-graded aluminum foam. Mater. Sci. Eng. A 618, 496–499 (2014) 5. L. Huang, H. Wang, D. Yang, F. Ye, Z.P. Lu, Effects of scandium additions on mechanical properties of cellular Al-based foams. Intermetallics 28, 71–76 (2012)

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6. X. Xia, H. Feng, X. Zhang, W. Zhao, The compressive properties of closed-cell aluminum foams with different Mn additions. Mater. Des. 51, 797–802 (2013) 7. J. Zhou, Z. Gao, A.M. Cuitino, W.O. Soboyejo, Effects of heat treatment on the compressive deformation behavior of open cell aluminum foams. Mater. Sci. Eng. A 386(1–2), 118–128 (2004) 8. X. Xia, W. Zhao, X. Feng, H. Feng, X. Zhang, Effect of homogenizing heat treatment on the compressive properties of closed-cell Mg alloy foams. Mater. Des. 49, 19–24 (2013) 9. H.F. Cheng, F.S. Han, Compressive behavior and energy absorbing characteristic of open cell aluminum foam filled with silicate rubber. Scr. Mater. 49(6), 583–586 (2003) 10. Y.W. Kwon, R.E. Cooke, C. Park, Representative unit-cell models for open-cell metal foams with or without elastic filler. Mater. Sci. Eng. A 343(1), 63–70 (2003) 11. J. Yuan, X. Chen, W. Zhou, Y. Li, Study on quasi-static compressive properties of aluminum foam-epoxy resin composite structures. Compos. B Eng. 79, 301–310 (2015) 12. Y. Li, Y. Wei, L. Hou, C. Guo, S. Yang, Fabrication and compressive behaviour of an aluminium foam composite. J. Alloy. Compd. 649, 76–81 (2015) 13. Q. Li, G. Jiang, C. Wang, J. Dong, G. He, Mechanical degradation of porous titanium with entangled structure filled with biodegradable magnesium in Hanks’ solution. Mater. Sci. Eng. C 57, 349–354 (2015) 14. S.N. Esfahani, M.T. Andani, N.S. Moghaddam, R. Mirzaeifar, M. Elahinia, Independent tuning of stiffness and toughness of additively manufactured titanium-polymer composites: simulation, fabrication, and experimental studies. J. Mater. Process. Technol. 238, 22–29 (2016)

Study on Heat Treatment Process of Tungsten-Plated Diamond Zhen Zhang, Hong Guo, Zhongnan Xie and Ximin Zhang

Abstract This study investigates the effects of the heat treatment process including different temperature rising rates, holding time, and holding temperature on the bonding strength of the tungsten-plated diamond coating. The heat treatment process is optimized by various analytical methods. The results show that tungsten carbide (WC) with strong bonding strength and high thermal conductivity can be obtained by holding the temperature at 1150 °C for 15 min. When the temperature rising rate is high, the total carbide content is high, and the tungsten carbide is amorphous. As the temperature rising rate decreases, the tungsten carbide content gradually decreases to zero, and the diamond is not graphitized.

1 Introduction With the development of modern science and technology, electronic instruments are gradually becoming more concentrated and miniaturized, resulting in an exponential increase in the heat dissipation per unit area of electronic components [1]. In order to improve the heat dissipation efficiency, a new type of heat dissipation material is essential in present. Through recent research, the diamond–copper composite is one of the best choices [2, 3]. In diamond composites, interface thermal resistance is an important factor affecting the thermal conductivity of materials [4]. Metallization of diamond surface is one of the best ways to reduce the thermal resistance of the interface, and the suitable coating structure will ensure the high thermal conductivity Z. Zhang · H. Guo (B) · Z. Xie · X. Zhang National Engineering Research Center for Nonferrous Metals Composites, General Research Institute for Non-ferrous Metals, Beijing 100088, China e-mail: [email protected] Z. Zhang e-mail: [email protected] Z. Xie e-mail: [email protected] X. Zhang e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_46

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of the material [5]. Among the different metal plating, due to the high thermal conductivity of tungsten carbide, the magnetron sputtering tungsten plating technology is relatively mature and the tungsten surface is ideal for tungsten plating [6]. For the tungsten-plated diamond prepared by the magnetron sputtering method, the combination of the tungsten plating layer and the diamond is only a physical combination, and the heat cannot be effectively transferred, so that a subsequent heat treatment is required to form a strong high thermal conductivity tungsten carbide combination, while avoiding the formation of graphitization and low thermal conductivity of tungsten carbide to achieve its proper performance [7]. For the formation of carbides, temperature is the most critical factor, so the holding temperature is quite important, and the relevant reaction kinetic conditions such as temperature rising rate and holding time are also factors to be studied. This paper investigates the reaction history and products of tungsten-plated diamond under different heat treatment processes and evaluates their performance.

2 Experimental Content and Method 2.1 Experimental Materials and Instruments The size of the raw material diamond is about 30 µm and its purity over 99%. The tungsten-plated diamond used in this experiment is prepared by magnetron sputtering, and it has a plating thickness of 30 nm in size. The tungsten-plated diamond powder was heat treated by a different heat treatment process using a helium vacuum furnace at a vacuum of 10−3 Pa. After the heat treatment, the phase of the material was analyzed by XRD diffractometer. The microstructure of the tungsten-plated diamond powder was observed and studied by scanning electron microscopy and energy dispersive spectroscopy.

2.2 Experimental Procedure In this experiment, the main factors of temperature rising rate, holding temperature and holding time were controlled. The effects of different heat treatment processes on tungsten-plated diamond powder were analyzed. We have set up five experimental groups. The specific heating process is shown in Table 1 and Fig. 1. The samples obtained by different heat treatment processes were analyzed by XRD, and the contents of the carbide, graphite, elemental tungsten were determined to analyze the experimental results.

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Fig. 1 Schematic diagram of the heat treatment process Table 1 Heat treatment process Time (min)

Temperature (°C) Process 1

Process 2

Process 3

Process 4

Process 5

0

25

25

25

25

25

60

600

600

600

600

600

90

750

750

900

750

900

120

900

900

1150

900

1150

135

987.5

960

1150

995

1150

150

1075

1020

1090

1150

180

1250

1150

1250

195

1250

1150

1250

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1250

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(a)

(b)

(c)

(d)

Nano tungsten particles

Fig. 2 a Tungsten-plated diamond topography; b W element distribution on tungsten-plated the surface of diamond; c surface morphology of tungsten-plated diamond before heat treatment; d surface morphology of tungsten-plated diamond after heat treatment

3 Results and Discussion 3.1 Coating Morphology and Heat Treatment Process Figure 1a–d shows the SEM results of tungsten-plated diamonds and distribution of tungsten on the surface. It can be seen from the electron micrograph that there are many small particles distributed in the coating, the distribution is relatively uniform, and the particle size is nanometer, after the heat treatment, the surface morphology of the diamond surface changed significantly. When the coating amplified 5000 times, the surface of the diamond was flocculated [8]. Which is because the uncoated layer is mainly composed of elemental tungsten, and the elemental tungsten is converted into carbide after heat treatment thus, it changes the morphology of the phase [9] (Fig. 2). The tungsten plating obtained by magnetron sputtering is less in combination with diamond. The bond between the tungsten-plated diamond coating and the diamond is mainly derived from the carbide formed by the reaction between the two. The thermal conductivity of WC is 121 W/(m K), while the thermal conductivity of W2 C is only

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Fig. 3 XRD diffractions corresponding to different heat treatment processes

36 W/(m K), The graphitization of diamond is also unfavorable for the combination of the coating and the diamond matrix. Therefore, it is desirable to obtain a tungstenplated diamond in which all of the tungsten is converted into WC rather W2 C and graphite are present after heat treatment.

3.2 Effect of Holding Time on Tungsten-Plated Diamond Phase In Fig. 3, curve 1–5 is the XRD curve of process No. 1–5, No. 6 is an untreated tungsten-plated diamond curve. Overall, relative to the untreated diamonds, the five processes essentially removed the elemental tungsten in the tungsten-plated diamond by the reaction but differed in the amount of product formed and the amount of carbide. First, according to the study of the effect on holding time, it can be seen from the comparison of No. 1 and No. 4 processes that at 1250 °C, there will be obvious graphitization phenomenon no matter the temperature is kept for 30 or 15 min. However, by reducing the holding time, the amount of graphite produced can be effectively reduced. However, at the same time, when the holding time is longer, more WC is formed, and the W2 C carbide is reduced, indicating that the W2 C is first formed in the heat treatment process, and is gradually converted into WC at a high temperature, the extension of the holding time is undoubtedly beneficial to the formation of WC. Comparing the No. 3 and No. 5 processes, it can be seen that when the holding temperature is 1150 °C, the holding time has no obvious influence on the graphitization phenomenon. But for the formation of carbides, it still follows the

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same rules: the increase in holding time will greatly increase the content of tungsten carbide.

3.3 Effect of Holding Temperature on the Phase of Tungsten-Plated Diamond In the heat treatment process, the holding temperature of No. 1 and No. 4 is 1250 °C, while the other three are 1150 °C. Comparing the XRD results of the two holding temperatures, it can be seen that the graphitization of diamond is very obvious at high temperature. It shows that the graphitization of tungsten-plated diamond is still affected by temperature, and the protection of the coating cannot completely avoid graphitization. At the same time, combined with literature research, the graphitization temperature of pure diamond under anaerobic conditions should be above 1500 °C [10], the added tungsten element is easily reacted with the diamond, so that some of the carbon element is detached from the diamond surface into the interior of the tungsten carbide lattice [11]. Corresponding to reducing the bonding force between carbon atoms, thereby reducing the barrier of diamond graphitization, that is, reducing the initial temperature of diamond graphitization. Therefore, in the process of heat treatment, the holding temperature should be lowered correspondingly on the basis of pure diamond data to avoid the influence of graphitization on material properties.

3.4 Effect of Temperature Rising Rate on Tungsten-Plated Diamond Phase Since the temperature rising rates of No. 1 and No. 4, No. 3 and No. 5 in the process are the same, thus select one of each set as a control, and the temperature rising rates of No. 5, No. 4, and No. 2 are successively decreased. From the XRD curve, first, when it is compared with No. 5 and No. 4, the relative content of tungsten carbide is higher, the graphitization is less obvious, and the peak area corresponding to tungsten carbide is relatively close. However, it can be seen that the peak width of the tungsten carbide of No. 5 is much larger than that of No. 4, indicating that the tungsten carbide formed during the heat treatment is in an amorphous state due to the faster reaction rate when the temperature rising rate is faster, at the same time, it may cause a certain thermal stress. The different content of tungsten carbide indicates that the amorphous structure of tungsten carbide formed in the early stage catalyzes the conversion of tungsten carbide to tungsten carbide when the temperature rises rapidly, so that better thermodynamic/kinetic conditions are obtained. Compared with the No. 4 and No. 2 processes, it can be seen that when the temperature rising rate is further reduced, the No. 2 process can almost completely

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avoid the formation of tungsten carbide. According to the literature, the conversion of the tungsten carbide is not only performed at 1100 °C or higher [12]. Indeed, there is a certain balance of conversion reaction at the same time as the formation of ditungsten carbide at 800 °C. As the temperature rising rate is reduced, the carbide conversion time is correspondingly increased, so that the most ideal tungsten-plated diamond completely free of tungsten carbide is obtained.

4 Conclusion 1. The tungsten element in tungsten-plated diamond reduces the temperature of diamond graphitization to 1250 °C, and the complete reaction of tungsten to tungsten carbide needs to be at the temperature above 1100 °C. Therefore, it is a reasonable process to keep the temperature at 1150 °C. 2. The optimal tungsten heat treatment process is as follow: The temperature rises from room temperature to 600 °C at 10 °C/min, then from 600 to 900 °C at 5 °C/min, then subsequently from 900 to 1150 °C at 4 °C/min and finally at 1150 °C for 15 min. When the temperature rising rate is high, the content of tungsten carbide and tungsten carbide is high, but the tungsten carbide is amorphous. As the temperature rising rate decreases, the tungsten carbide content gradually decreases to zero with a complete conversion of elemental tungsten to tungsten carbide in the coating without the diamond graphitized. Acknowledgement The authors would like to acknowledge the financial supports from the Ministry of Science & Technology of China (the National Key Research and Development Program of china No. 2017YFB0406202).

References 1. C. Zweben, Thermal materials solve power electronics challenges. Power Electron. Technol. 32(2), 40–47 (2006) 2. S.Q. Liu, X.M. Liu, Z.Q. Zhang et al., Study of electroless nickel plating on micro-electronics packaging and its applications. Electroplat. Finish. 20(3) (2005) 3. Q.L. Shang, J.M. Tao, M.-C. Xu et al., Research progress of diamond-Cu composite material for electronic packaging. Electron. Process Technol. (5), 56–61 (2009) 4. A.M. Abyzov, M.J. Kruszewski, Ł. Ciupi´nski et al., Diamond–tungsten based coating–copper composites with high thermal conductivity produced by pulse plasma sintering. Mater. Des. 76, 97–109 (2015) 5. J. Li, H. Zhang, Y. Zhang et al., Microstructure and thermal conductivity of Cu/diamond composites with Ti-coated diamond particles produced by gas pressure infiltration. J. Alloy. Compd. 647, 941–946 (2015) 6. J. Li, H. Zhang et al., On the thermal conductivity of Cu/diamond composite of diamond particles with tungsten coating. J. Funct. Mater. 47(1), 1034–1037 (2016)

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7. Q. Kang, X. He, S. Ren et al., Microstructure and thermal properties of copper–diamond composites with tungsten carbide coating on diamond particles. Mater. Charact. 105, 18–23 (2015) 8. A.M. Abyzov, S.V. Kidalov, F.M. Shakhov, High thermal conductivity composites consisting of diamond filler with tungsten coating and copper (silver) matrix. J. Mater. Sci. 46(5), 1424–1438 (2011) 9. C. Zhang, R. Wang, C.Q. Peng et al., Effect of surface tungsten plating on thermal conductivity of diamond/copper composites. Rare Met. Mater. Eng. 45(10), 2692–2696 (2016) 10. Z.J. Qiao, Nanodiamonds Graphitization Transformation and Preparation and Properties of Nano-Diamond/Copper Composites (Tianjin University, 2007) 11. H.T. Weigh, Y.X. Wu, J.H. Chen et al., Study on preparation process and antioxidation performance of vacuum evaporated tungsten diamond. Superhard Mater. Eng. (2) (2018) 12. H.P. Lkixev, Metal Binary Phase Diagram Manual (Chemical Industry Press, 2009)

Effect of Sc Modification and Pulping Process on Semi-Solid Structure of A356 Aluminum Alloy Yuxin Zhang, Hengbin Liao, Yong Dong, Anfu Chen, Xiaoling Fu and Zhengrong Zhang

Abstract In this paper, A356 aluminum alloy was used as the research object, and the traditional mixing method was used to explore the parameters of stirring pulping process and the effect of Sc metamorphism on semi-solid slurry during continuous cooling. The experimental results showed that with the increase of shear rate, more broken dendrites were obtained, resulting in smaller equivalent diameter of grain size and higher shape factor (roundness) of semi-solid slurry. In this experiment, the best effect was obtained while the shear rate was 70 s−1 . For the initial stirring temperature, stirring was started at 630 °C and decreased to the optimal temperature 595 °C at which the slurry was obtained. Besides, the temperature-dropping velocity had a significant effect on the slurry quality. When the decrease of temperature (the best effect at 1.94 °C/min) was slower and the shear time was more fully, the dendrite precipitation grew slower. Therefore, the dendrites that had just crystallized out had been sheared and broken in the future, resulting in a near-spherical shape. When the number of crystal grains (nucleation nuclei) was larger, the crystal grain equivalent diameter of the slurry was smaller and the shape factor was larger. In terms of the semi-solid slurry A356 containing Sc element, it was found that the grain size was smaller and the degree of roundness was higher.

1 Introduction The quality of the semi-solid slurry determines the performance of the part. The evaluation criteria of the semi-solid slurry are the grain size of the semi-solid slurry, the roundness of the grain, the pores, and slag inclusions inside the slurry melt. Therefore, one of the keys to semi-solid forming technology is to prepare an excellent metal semi-solid slurry [1]. In this experiment, the commercial aluminum alloy A356 Y. Zhang (B) · H. Liao · Y. Dong · A. Chen · X. Fu · Z. Zhang School of Materials and Energy, Guangdong University of Technology, Guangzhou 510006, Guangdong, China e-mail: [email protected] Z. Zhang e-mail: [email protected] © Springer Nature Singapore Pte Ltd. 2019 Y. Han (ed.), Physics and Engineering of Metallic Materials, Springer Proceedings in Physics 217, https://doi.org/10.1007/978-981-13-5944-6_47

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and Sc refined and modified A356 was used as the object, and the traditional mixing method was used to explore the stirring pulping process parameters and the influence of Sc element on the semi-solid slurry in the continuous cooling process [2–8].

2 Results and Discussion 2.1 Pulp Experiment Process The semi-solid mechanical pulping device used in this experiment is shown in Fig. 1. The experimental process is as follows: the bulk A356 aluminum ingot Al–20Si master alloy had been preheated to 250 °C clay graphite crucible and melted to 740–75 °C in well type crucible resistance furnace. At the beginning of melting at 620–630 °C, we sprayed a layer of dried covering on the solution surface to avoid inhaling too much oxygen and hydrogen. When the temperature reaches 700 °C, we scraped the slag and press the aluminum foil around the dried Al–50Mg master alloy into the bell hood. The alloy solution was made to be homogenized by slight rotation and stirring, and then a layer of covering agent was sprayed. After 5 min of rest, the slag was removed, and the graphite rotor was inserted into the alloy melt and passed into the high purity argon for impurity removal refining. Then the covering agent was sprayed 5 min later. Continue to heat to 740–750 °C, we removed the surface covering agent, and pressed the prepared Al–2Sc master alloy into the melt of the alloy. After the melt was completely melted, we mixed evenly with a mixing rod and sprinkle with a coating of covering agent for heat preservation. Let it sit for five minutes, and we removed the slag. We inserted the graphite rotor into the alloy melt and passed into the high purity argon for the second impurity refining. After 5 min, we spread the covering agent. After 5 min of statics, the melt temperature was controlled at 710–720 °C, and the agitator preheated to 300–350 °C was inserted into the alloy solution and stirred at a certain shear rate until the appropriate initial stirring temperature was reached. At 595 °C, the quenched water was removed and the quenched sample was ground from 600# to 2000# on the metallographic sandpaper. After polishing with diamond grinding paste (W0.5), the quenched sample was corroded by Kohler reagent. The technological parameters designed in this experiment are affected by the addition of SC elements, as shown in Table 1. In addition, the equivalent diameters and shape factors are obtained by calculating the 150–200 grains in the metallographic photographs of each experimental group. We, respectively, express the near-spherical grain size and the roundness of the obtained grains. The microstructure of quenched samples is calculated by ImagePro Plus software and based on the formula (1), (2). The grain equivalent diameter and shape factor of the slurry is obtained. Equivalent diameter : D 

i  Ai n2 π n

(1)

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Fig. 1 Mechanical stirring pulping device Table 1 Experiment process parameter Group No.

Process parameters, Sc content Shear rate (s−1 )

Initial stirring temperature (°C)

Temperature drop rate (°C/s)

Sc content (wt%)

A1

24

630

3.33