Applied Science, Faculty of
Engineering, School of (Okanagan)
DSpace
UBCO
Gautam, Sadhak
Master of Applied Science - MASc
University of British Columbia
Hadfield steels, also known as high manganese (Mn) steels (HMS), are often used to fabricate components for crushing equipment, because of HMS’s high wear and impact resistance arising from the alloy’s high hardness and toughness. The composition of Hadfield steels has a significant effect on mechanical properties of the alloys. In this research, the effect of varying Mn (12-23 wt. %) on the ternary Fe-Mn-C system, microstructure and mechanical properties of HMS alloys was studied. Liquidus projection of ternary Fe-Mn-C diagram revealed the decrease in liquidus temperature with an increase in Mn from 12 to 23 wt. %. Microstructure analysis revealed that an alloy with 12 wt. % Mn had intermittent carbides at the grain boundaries. For alloys with 18 and 23 wt.% Mn, discontinuous carbides had a lamellar structure and formed at the grain boundaries and intergranular regions. The average grain size increased with an increase in the Mn concentration in the alloys. Energy dispersive spectroscopy (EDS) map, line, and point/area scans revealed the accumulation of chromium (Cr) and Mn on the carbide sites in all the alloys. (Fe,Cr)₂₃C₆ intermetallic compound was present in all the alloys, with a slight variation of Cr and Fe in localized areas. The Fe₁․₈Mn₁․₂C was present in alloys with 12 and 18 wt. % Mn, whereas Fe₁․₁Mn₃․₉C₂ was present in an alloy with 23 wt. % Mn. Mechanical testing (impact and tensile) revealed that the Charpy-impact values reached the highest value at 23 wt. % Mn addition and were lower for other alloy compositions. With an increase in Mn content in the alloys, the yield strength increased, whereas the ultimate tensile strength and elongation decreased. The hardness increased with an increase in Mn wt. % in the alloys. Finally, the weight loss of material during pseudo-wear testing decreased with an increase of Mn from 12 to 18 wt.% then remained constant for an alloy with 23 wt. % Mn.
ii INVESTIGATING THE EFFECT OF Mn CONTENT ON THE MICROSTRUCTURE AND MECHANICAL PROPERTIES OF HIGH MANGANESE STEEL by Sadhak Gautam Bachelor of Technology, SRM University, 2017 A THESIS SUBMITTED IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF MASTER OF APPLIED SCIENCE in THE COLLEGE OF GRADUATE STUDIES (Mechanical Engineering) THE UNIVERSITY OF BRITISH COLUMBIA (Okanagan) March 2022 © Sadhak Gautam, 2022 ii The undersigned certify that they have read, and recommend to the College of Graduate Studies for acceptance, a thesis entitled: INVESTIGATING THE EFFECT OF Mn CONTENT ON THE MICROSTRUCTURE AND MECHANICAL PROPERTIES OF HIGH MANGANESE STEEL submitted by Sadhak Gautam in partial fulfillment of the requirements of the degree of Master of Applied Science. Dr. Lukas Bichler, School of Engineering Supervisor, Professor Dr. Abbas Milani, School of Engineering Supervisory Committee Member, Professor Dr. Dimitry Sediako, School of Engineering Supervisory Committee Member, Professor Dr. Stephen O’Leary, School of Engineering University Examiner, Professor iii Abstract Hadfield steels, also known as high manganese (Mn) steels (HMS), are often used to fabricate components for crushing equipment, because of HMS’s high wear and impact resistance arising from the alloy’s high hardness and toughness. The composition of Hadfield steels has a significant effect on mechanical properties of the alloys. In this research, the effect of varying Mn (12-23 wt. %) on the ternary Fe-Mn-C system, microstructure and mechanical properties of HMS alloys was studied. Liquidus projection of ternary Fe-Mn-C diagram revealed the decrease in liquidus temperature with an increase in Mn from 12 to 23 wt. %. Microstructure analysis revealed that an alloy with 12 wt. % Mn had intermittent carbides at the grain boundaries. For alloys with 18 and 23 wt.% Mn, discontinuous carbides had a lamellar structure and formed at the grain boundaries and intergranular regions. The average grain size increased with an increase in the Mn concentration in the alloys. Energy dispersive spectroscopy (EDS) map, line, and point/area scans revealed the accumulation of chromium (Cr) and Mn on the carbide sites in all the alloys. (Fe,Cr)23C6 intermetallic compound was present in all the alloys, with a slight variation of Cr and Fe in localized areas. The Fe1.8Mn1.2C was present in alloys with 12 and 18 wt. % Mn, whereas Fe1.1Mn3.9C2 was present in an alloy with 23 wt. % Mn. Mechanical testing (impact and tensile) revealed that the Charpy-impact values reached the highest value at 23 wt. % Mn addition and were lower for other alloy compositions. With an increase in Mn content in the alloys, the yield strength increased, whereas the ultimate tensile strength and elongation decreased. The hardness increased with an increase in Mn wt. % in iv the alloys. Finally, the weight loss of material during pseudo-wear testing decreased with an increase of Mn from 12 to 18 wt.% then remained constant for an alloy with 23 wt. % Mn. v Lay Summary Mining is a process of extracting materials such as coal, metals, etc. from rocks. An explosive process is used for breaking the native rocks, followed by a crushing process, to reduce the rock’s size for mineral extraction. During the different stages of crushing, the rock size reduces from ~ 1.2 m to ~ 0.5 cm using impact and grinding processes. When the rocks impact the crusher’s surface, wear (loss of material) occurs. This leads to the degradation of the impact crusher’s performance, followed by the necessary shut down and mandatory replacement of the crusher’s parts. Therefore, this thesis focused on studying the effect of manganese (Mn) on the mechanical properties of Hadfield steel materials used in rock crushers. Mechanical tests were carried out to evaluate the performance of new alloys, followed by detailed material characterization of all alloy alloys. vi Preface This thesis on investigating the effect of Mn content in Hadfield steel was written in a partial requirement for the fulfillment of a master’s degree (MASc) at the University of British Columbia (UBC), Okanagan. The optical microscopy analysis, pseudo-wear test, and hardness test were performed by Sadhak Gautam in Dr. Lukas Bichler’s research laboratory at the University of British Columbia, Okanagan. The SEM/EDS analysis was carried out at the FILTER laboratory at UBC, Okanagan. X-ray Diffraction scans were carried out by Anita Lam at the Department of Chemistry, UBC, Vancouver, and phases were analyzed by Sadhak Gautam using the X'Pert High Score software available at Dr. Lukas Bichler’s laboratory. Tensile testing was performed at Composites Research Network at UBC, Okanagan by Sadhak Gautam. Charpy-impact testing was performed by Sadhak Gautam at UBC, Kelowna, BC, Canada under the supervision of the School of Engineering Technical Support team. The manuscripts were written and edited by Sadhak Gautam under the supervision of Dr. Lukas Bichler. vii Table of Contents Abstract ......................................................................................................................... iii Lay Summary ................................................................................................................. v Preface ........................................................................................................................... vi Table of Contents ......................................................................................................... vii List of Tables ................................................................................................................. xi List of Figures .............................................................................................................. xiii List of Abbreviations and Symbols ........................................................................... xviii List of Elements and Compounds ............................................................................... xix Acknowledgment .......................................................................................................... xx Dedication ................................................................................................................... xxii Chapter 1: Introduction ................................................................................................. 1 1.1 Background and Statement of Problem ................................................................. 1 1.2 Objective and Scope of Research .......................................................................... 2 Chapter 2: Literature Review ........................................................................................ 4 2.1 General Background ............................................................................................. 4 Mining Process ............................................................................................. 5 Types of Wear and Required Properties of Wear Resistant Material .............. 7 Material Options for Crushing Applications .................................................. 8 2.2 Introduction to Hadfield Steels.............................................................................. 9 viii 2.3 Chemical Composition........................................................................................ 10 Effect of Manganese Content ...................................................................... 11 Effect of Carbon Content ............................................................................ 15 Effect of Other Alloying Elements .............................................................. 17 2.4 Casting Parameters ............................................................................................. 21 Effect of Pouring Temperature .................................................................... 21 2.5 Heat-treatment of Hadfield Steels ....................................................................... 23 2.6 Work-hardening of Hadfield Steels ..................................................................... 25 2.7 Chapter Summary ............................................................................................... 27 Chapter 3: Experimental Procedure ........................................................................... 28 3.1 Casting and Heat-treatment Process .................................................................... 28 3.2 Microstructure Analysis ...................................................................................... 30 Sample Preparation ..................................................................................... 30 Optical Microscopy ..................................................................................... 31 Scanning Electron Microscopy .................................................................... 31 3.3 Phase Analysis .................................................................................................... 32 X-Ray Diffraction ....................................................................................... 32 3.4 Mechanical Property Evaluation ......................................................................... 33 Charpy-Impact Testing ................................................................................ 33 Tensile Testing............................................................................................ 34 Pseudo-Wear Testing .................................................................................. 35 Microhardness Testing ................................................................................ 37 ix Chapter 4: Results and Discussion .............................................................................. 38 4.1 Fe-Mn-C System Analysis .................................................................................. 38 Liquidus Projection ..................................................................................... 38 Phase analysis ............................................................................................. 40 4.2 Microstructure Analysis ...................................................................................... 47 Optical Microscopy ..................................................................................... 47 Scanning Electron Microscopy - General Microstructure............................. 55 X-ray Energy Dispersive Spectroscopy ....................................................... 58 4.3 Phase Analysis .................................................................................................... 66 4.4 Mechanical Testing ............................................................................................. 68 Charpy-impact Testing ................................................................................ 68 Tensile Testing............................................................................................ 70 Microhardness............................................................................................. 71 Pseudo-wear Testing ................................................................................... 73 Chapter 5: Conclusions and Recommended Future Work ......................................... 74 5.1 Fe-Mn-C System, Microstructure, and Phase Analysis ........................................ 74 5.2 Mechanical Properties ......................................................................................... 75 5.3 Recommendations for Future Work .................................................................... 75 Low-Temperature Charpy-impact and Tensile Testing ................................ 75 Transmission Electron Microscope analysis ................................................ 76 Impact Wear Testing ................................................................................... 76 Bibliography ................................................................................................................. 77 x Appendices .................................................................................................................... 84 Appendix A ................................................................................................................ 84 Appendix B ................................................................................................................ 86 xi List of Tables Table 2.1: Standard ASTM grades of Hadfield steels [4], [22]. ........................................... 11 Table 2.2: Chemical composition of alloys with 12 and 16 wt. % Mn [6]. .......................... 13 Table 2.3: Performance characteristics of 12 wt.% Mn and 16 wt.% Mn steels [6].............. 14 Table 2.4: Dependence of alloy grain size on the solidus and liquidus temperatures [30]. ... 17 Table 2.5: Chemical composition of 12 wt. % Mn steels with and without Cr [28]. ............ 18 Table 3.1: Chemical composition of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn. .......................................................................................................................................... 28 Table 3.2: XRD test parameters. ......................................................................................... 32 Table 3.3: Stages of pseudo-wear test. ................................................................................ 37 Table 4.1: Liquidus temperature and superheat of Fe-12.14Mn-1.05C, Fe-17.7Mn-1.13C, and Fe-22.78Mn-1.15C alloys. .................................................................................................. 40 Table 4.2: Volume percentage of cementite in Fe-12.14Mn-1.05C, Fe-17.7Mn-1.13C, and Fe-22.78Mn-1.15C at 650℃, 750℃, and 900℃. ..................................................................... 41 Table 4.3: Wt.% of C and Mn in alloy with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn addition. ............................................................................................................................. 51 Table 4.4 Average Grain Size in alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn addition. ............................................................................................................................. 52 Table 4.5: Phase area fraction in alloys with 12 wt. %, 18 wt. %, and 23 wt. % Mn addition. .......................................................................................................................................... 53 Table 4.6: Chemical composition (wt. %) in an alloy with 12 wt. % Mn. ............................ 65 Table 4.7: Chemical composition (wt. % ) in an alloy with 18 wt. % Mn. ........................... 65 xii Table 4.8: Chemical composition (wt. % ) in an alloy with 23 wt. % Mn. ........................... 66 Table 4.9: Summary of microstructure analysis and wt.% of Mn and C in alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn addition. ................................................................ 69 xiii List of Figures Figure 1.1: Flowchart illustrating the scope of the present research. ..................................... 3 Figure 2.1: Mining contribution index (MCI-Wr) score by country 2016 [9]. ....................... 4 Figure 2.2: Stages of crushing process [1]. ........................................................................... 5 Figure 2.3: Working principle of: (A) jaw crusher, and (B) gyratory crusher [1]. .................. 6 Figure 2.4: Working principle of: (A) impact crusher, and (B) roller crusher [1]................... 6 Figure 2.5: Two types of abrasion wear: (A) two-body abrasion, (B) three-body abrasion [12]. ............................................................................................................................................ 7 Figure 2.6: Picture of: (A) truck bed is abraded by the free flow of dumped ore or minerals under relatively low wearing stress, and (B) wear of roller crusher under high stress [12]..... 8 Figure 2.7: Toughness vs. abrasion resistance [14], [15]. ...................................................... 9 Figure 2.8: Jaw crusher plate made of Hadfield steel [17]. .................................................. 10 Figure 2.9: Fe-Mn binary phase diagram [23]. .................................................................... 12 Figure 2.10: Effect of Mn content on the mechanical properties and microstructure of Hadfield steels [18]. .......................................................................................................................... 13 Figure 2.11: Micrograph of alloys with: (A) 12 wt. %Mn addition, and (B) 16 wt. % Mn addition [6]. ....................................................................................................................... 14 Figure 2.12: Effect of varying carbon (wt. %) on the mechanical properties and microstructure of Hadfield steels [18]. ....................................................................................................... 15 Figure 2.13: Phase diagram of Fe-C-Mn [29]. .................................................................... 16 Figure 2.14: Effect of varying chromium (wt.%) on the mechanical properties and microstructure of Hadfield steels [18]. ............................................................................... 18 xiv Figure 2.15: Variation in grain size (µm) and hardness (HV) for 11.6 wt. % Mn alloy with and without Cr [28]................................................................................................................... 19 Figure 2.16: Microstructure of 12% Mn steel poured at 1550℃: (A) edge, (B) middle, and (C) mounting [25]. ................................................................................................................... 22 Figure 2.17 Microstructure of 12% Mn steel poured at 1450℃: (A) edge, (B) middle, and (C) mounting [25]. ................................................................................................................... 22 Figure 2.18: Microstructure of 12% Mn steel poured at 1380 ℃: (A) edge, (B) middle, and (C) mounting [25]. ............................................................................................................. 22 Figure 2.19: Effect of pouring temperature on the degree of segregation of Mn, Cr and Si [25]. .......................................................................................................................................... 23 Figure 2.20: Optical micrographs of ASTM A128 grade B-3 steel at 500X magnification : (A) before heat-treatment, and (B) after heat-treatment at 1120 ℃ [4]. ..................................... 24 Figure 2.21: Microstructure of cast Hadfield steel after solution treatment followed by water quenching with water maintained at: (A) 1200 ℃ with 80 minutes of hold time, and (B) 1150 ℃ with 240 minutes of hold time [36]. ............................................................................... 25 Figure 2.22: Variation in hardness (HV) vs. shot peening duration [44]. ............................. 27 Figure 3.1: Pictures of casting block: (A) front-view, and (B) top-view (all dimensions in cm). .......................................................................................................................................... 29 Figure 3.2: Heat-treatment chart of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn. .......................................................................................................................................... 29 Figure 3.3: Picture of cylindrical microscopy sample. ........................................................ 30 Figure 3.4: Field emission scanning electron microscopy setup. ......................................... 31 Figure 3.5: V-notch Charpy-impact sample size [46]. ......................................................... 33 xv Figure 3.6: Charpy-impact tester. ....................................................................................... 34 Figure 3.7: Picture of: (A) sub-size tensile sample dimensions [47], and (B) prepared tensile sample. ............................................................................................................................... 35 Figure 3.8: Picture of: (A) Struers Tegramin-30 auto-polisher, and (B) samples mounted on auto-polisher. ..................................................................................................................... 36 Figure 3.9: Sample holder with position P1 to P6. .............................................................. 36 Figure 3.10: Microhardness tester. ...................................................................................... 37 Figure 4.1: Liquidus projection of Fe-Mn-C system (all temperatures in Kelvin). ............... 39 Figure 4.2: Ternary phase diagram of Fe-Mn-C system at 900 °C. ...................................... 41 Figure 4.3: Ternary phase diagram of Fe-Mn-C system at 900 °C (Mn: 0-50 mass %). ....... 42 Figure 4.4: Ternary phase diagram of Fe-Mn-C system at 750 °C. ...................................... 42 Figure 4.5: Ternary phase diagram of Fe-Mn-C system at 750 °C (Mn: 0-50 mass %). ....... 43 Figure 4.6: Ternary phase diagram of Fe-Mn-C system at 650 °C. ...................................... 43 Figure 4.7: Ternary phase diagram of Fe-Mn-C system at 650 °C (Mn: 0-50 mass %). ....... 44 Figure 4.8: Ternary phase diagram of Fe-Mn-C system at 500 °C. ...................................... 45 Figure 4.9: Ternary phase diagram of Fe-Mn-C system at 500 °C (Mn: 0-50 mass %). ....... 46 Figure 4.10: Ternary phase diagram of Fe-Mn-C system at 250 °C. .................................... 46 Figure 4.11: Ternary phase diagram of Fe-Mn-C system at 250 °C (Mn: 0-50 mass %). ..... 47 Figure 4.12: Optical micrographs of an alloy with 12 wt. % Mn at: (A) 50X, (B) 100X, (C) 500X, and (D) 500X magnification. ................................................................................... 48 Figure 4.13 Optical micrographs of an alloy with 18 wt. % Mn at: (A) 50X, (B) 100X, and (C) 500X magnification. ........................................................................................................... 49 xvi Figure 4.14: Optical micrographs of an alloy with 23 wt. % Mn at: (A) 50X, (B) 100X, (C) 500X, (D) 500X , (E) 500X, and (F) 500X magnification. .................................................. 50 Figure 4.15: Optical micrographs of an alloy 12 wt. % Mn: (A) without color threshold, and (B) with a color threshold. .................................................................................................. 53 Figure 4.16: Optical micrographs of an alloy with 18 wt. % Mn: (A) without color threshold, and (B) with a color threshold. ........................................................................................... 53 Figure 4.17: Optical micrographs of an alloy with 23 wt. % Mn: (A) without color threshold, and (B) with a color threshold. ........................................................................................... 54 Figure 4.18: Scanning electron micrographs of an alloy with 12 wt. % Mn at: (A) 500X, (B) 1000X, and (C) 500X magnification. .................................................................................. 55 Figure 4.19: Scanning electron micrographs of an alloy with 18 wt. % Mn at: (A) 500X, (B) 1000X, (C) 500X, and (D) 500X magnification. ................................................................. 56 Figure 4.20: Scanning electron micrographs of an alloy with 23 wt. % Mn at: (A) 500X, (B) 2000X, and (C) 1000X magnification. ................................................................................ 57 Figure 4.21: EDS map scan of an alloy with 12 wt. % Mn. ................................................. 59 Figure 4.22: EDS map scan of an alloy with 18 wt. % Mn. ................................................. 60 Figure 4.23: EDS map scan of an alloy with 23 wt. % Mn. ................................................. 61 Figure 4.24: EDS line scan of an alloy with 12 wt. % Mn. .................................................. 62 Figure 4.25: EDS line scan of an alloy with 18 wt. % Mn. .................................................. 63 Figure 4.26: EDS line scan of an alloy with 23 wt. % Mn. .................................................. 64 Figure 4.27: EDS point/area scan of an alloy with 12 wt. % Mn. ........................................ 65 Figure 4.28: EDS point/area scan of an alloy with 18 wt. % Mn. ........................................ 65 Figure 4.29: EDS point/area scan of an alloy with 23 wt. % Mn. ........................................ 66 xvii Figure 4.30: XRD pattern of an alloy with 12 wt. % Mn. .................................................... 67 Figure 4.31: XRD pattern of an alloy with 18 wt. % Mn. .................................................... 67 Figure 4.32: XRD pattern of an alloy with 23 wt. % Mn. .................................................... 68 Figure 4.33: Charpy-impact values of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn. .................................................................................................................................... 69 Figure 4.34: Tensile testing results of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn. .................................................................................................................................... 70 Figure 4.35: Average hardness of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn. .......................................................................................................................................... 71 Figure 4.36: Average hardness value of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn during the pseudo-wear test. ......................................................................................... 72 Figure 4.37: Weight loss percentage of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn Vs time. ....................................................................................................................... 73 Figure A.1: Optical micrographs of an alloy with 12 wt. % Mn at: (A) 50X, (B)100X, (C)500X, and (D) 500X magnification……………………………………………………………........84 Figure A.2: Optical micrographs of an alloy with 18 wt. % Mn at: (A) 50X, (B)100X, (C)500X, and (D) 500X magnification………………………………………………………………....84 Figure A.3: Optical micrographs of an alloy with 23 wt. % Mn at: (A) 50X, (B)100X, (C)500X, and (D) 500X magnification……………………………………………………….………...85 Figure B.1: EDS map scan of an alloy with 12 wt. % Mn……………………………………86 Figure B.2: EDS map scan of an alloy with 18 wt. % Mn……………………………………87 Figure B.3: EDS map scan of an alloy with 23 wt. % Mn……………………………………88 xviii List of Abbreviations and Symbols ASTM American society for testing materials EDS Energy-dispersive x-ray spectroscopy FE Field emission HMS High manganese steel hr Hour MCI Mining contribution index RPM Rotation per minute SEM Scanning electron microscopy SHT Solution heat-treatment XRD X-ray diffraction cm Centimetre J Joule m Meter min Minute ml/sec Milliliter per second mm Millimeter MPa Megapascal N Newton sec Second Vs. Versus Wt. % Weight percentage ɣ Austenitic Fe α Ferritic Fe xix List of Elements and Compounds C Carbon N Nitrogen Al Aluminum Si Silicon P Phosphorus S Sulphur Cr Chromium Mn Manganese Fe Iron Ni Nickel Cu Copper Nb Niobium Mo Molybdenum SiC Silicon carbide xx Acknowledgment I would like to express my great appreciation for the all support I received from my supervisor, Dr. Lukas Bichler throughout my master’s. I will always be grateful to him for giving me this wonderful opportunity to grow as a student, teacher, and researcher. Even when the offices were closed, he was always virtually around to support and guide me in these unprecedented times. I am sure no other supervisor would have been able to support their students to this extent. I would like to express deep gratitude to Unicast, Kelowna for sponsoring this research work. I would like to thank Mr. Abhinav Karanam of Unicast, Kelowna, for his continuous support in this research work. He always helped and guided me on each step of experiment and manuscript work. This thesis would not be possible without his great help and support. I would like to thank Dr. Dimitry Sediako for giving me access to the Thermo-calc software. Special thanks to Dr. Dimitry Sediako and Dr. Joshua Stroh for guiding me with the Thermo-Calc calculation. I would like to thank each member of Dr. Lukas Bichler’s research group for all the fruitful discussions that inspired many new ideas for this research work. I sincerely thank Dr. Anil Prasad for training me on the experimental work and sharing his knowledge and experience. Thanks a lot to Somi and Lava for training me on X’Pert high score software. Special thanks to Dr. Abbas Milani for permitting me to use the tensile testing machine available at Composites Research Network, UBC Okanagan. I am thankful to Dr. Kristian Mackowiak and Somi Doja for proofreading the manuscript. xxi I would like to thank each member of the UBC technical support team for all the support and guidance. I thank Bora Acun and Kim Nordstrom for supervising the tensile testing and Charpy-impact testing. I would like to thank Sina Kianfar for sharing the extensometer attachments for tensile testing. Thanks to Anita Lam from the UBC Vancouver Department of Chemistry for running my XRD scans, and thanks to Dr. Sudip Shrestha for his guidance on SEM/EDS. xxii Dedication To my family. 1 Chapter 1: Introduction This chapter provides a general introduction to the mining sector and current materials used in crushing equipment. Then, the framework for this research work’s objective and scope is outlined. 1.1 Background and Statement of Problem Breaking of rocks to extract minerals from them is one of the biggest challenges in the mining industry. Explosive processes are often used for breaking large rocks, with the subsequent processing by non-explosive methods [1]. “Crushing” is a non-explosive method for sizing rocks to enable mineral extraction [2]. During crushing, the rocks impact or slide on the crusher’s teeth, liners, cones, etc., and subsequently fracture. The impact, however, results in the wearing of the crusher’s parts. Wear is the most common reason for the frequent replacement of the crusher’s parts and results in a costly shutdown of operations. As a result, a suitable selection of the crusher’s material is essential to increase productivity and reduce downtime [3]. Recent developments in specialty alloy metallurgy yielded various materials for use in rock crushing applications, such as high chromium iron alloys and Hadfield steels. High chromium iron alloys show excellent resistance to wear, but are expensive [2], [4], and crack during repetitive impact loading [5]. Hadfield steels, also known as austenitic/high manganese steels, are commonly used in mining applications due to their good resistance against wear, high toughness, and high ductility [2]–[4], [6]–[8]. Under impact loading, Hadfield steel undergoes work hardening without becoming brittle, thus being suitable for impact applications [2]. Frequent replacement of the crusher’s teeth or liners can be avoided by improving the 2 mechanical properties of the Hadfield steels, which can result in improved rock crushing productivity. Therefore, for the economic benefit of the industry, it is necessary to work on further optimizing the mechanical properties of Hadfield steels. 1.2 Objective and Scope of Research This research focused on exploring the effect of varying the Manganese (Mn) content on the Fe-Mn-C system, microstructure, and mechanical properties of Hadfield steels. Castings were prepared with Mn content of 12, 18, and 23 wt.%. In this analysis of the ternary Fe-Mn-C system, the effect of Mn on the liquidus temperature and phase evolution of solid phases was studied using the liquidus and phase diagram created in Thermo-Calc software, respectively. The effect of Mn on the microstructure was investigated using optical microscopy and scanning electron microscopy (SEM). Elements and phases were analyzed using the energy-dispersive x-ray spectroscopy (EDS) and x-ray diffraction (XRD), respectively. Lastly, the effect of Mn addition on the mechanical properties such as tensile, hardness, wear, and high-strain toughness (Charpy-impact) was studied. Figure 1.1 shows the overview of the research work done for this study. 3 Figure 1.1: Flowchart illustrating the scope of the present research. Effect of varying Managanse content in Hadfield SteelFe-Mn-C System AnalysisLiquidus ProjectionPhase DiagramMicrostruture AnalysisOptical MicroscopyGrain SizePhases Area FractionScanning Electron MicroscopyElemental and Phase AnalysisX-Ray Energy Dispersive Spectroscopy X-Ray DiffractionMechanical Properties EvaluationPseudo-Wear TestingTensile TestingMicrohardnessTestingCharpy-impact Testing4 Chapter 2: Literature Review This chapter provides a general background and introduction to the mining industry and its material requirements and options for crushing equipment. A detailed literature review on Hadfield steels, their casting parameters, heat-treatment temperatures, chemical composition, and their application in the mining industry are presented. 2.1 General Background Mining is one of the major industries for the economic development of many industrialized and developing nations alike. The Mining Contribution Index (MCI) tracks the contribution of mining industries to a country’s national economic development. Figure 2.1 shows the MCI of the world in 2016. The mining industry has played a significant role in establishing Canada at the international level as a resource-based economy. Mining industries contributed to the gross domestic product (GDP) of Canada by approximately 3.2% ($60.3 billion) in 2015 [9] and 5% ($109 billion) in 2018 [10]. Figure 2.1: Mining contribution index (MCI-Wr) score by country 2016 [9]. 5 Mining Process Mining involves breaking and processing rocks to extract desirable commodities such as coal or metals. During the mining process, when the hard rocks are handled by the crusher’s parts, wear occurs [3]. Breaking of rocks for the extraction is done in two stages using distinct methods: (1) Explosive and (2) Non-explosive methods [1]. Firstly, the rocks are broken down using the explosive method. The energy consumption in explosive techniques is relatively high because explosive tends to release sudden plastic energy [2]. The non-explosive methods are used for further reduction of rocks down to 0.5 cm using rock crushers, pulverizers, and hammers to assist the extraction process [1]. The crushing process is further classified into three stages- (1) Primary, (2) Secondary, and (3) Tertiary, as shown in Figure 2.2. Figure 2.2: Stages of crushing process [1]. 6 In the primary phase of a crushing process, the rocks of up to 1.5 m in size feed into the jaw crusher to reduce their size to 10-20 cm. The most common types of crushers are jaw and gyratory crushers, as shown in Figure 2.3 (A) and (B), respectively [1]. During the secondary crushing stage, the discharge of the primary stage is further reduced down to 0.5 to 2 cm. These crushers are smaller in size as compared to the primary stage crushers. The most common types of crushers used in this stage are roller crushers and impact crushers, as shown in Figure 2.4 [1]. A B Figure 2.3: Working principle of: (A) jaw crusher, and (B) gyratory crusher [1]. A B Figure 2.4: Working principle of: (A) impact crusher, and (B) roller crusher [1]. 7 In the last crushing stage, the discharge of the secondary crusher is analyzed. In case the size of discharge is too coarse, then the oversized discharge is sent back to the second stage for re-crushing. The remaining discharge is further processed using the small cone crushers and is sent for milling [1]. Types of Wear and Required Properties of Wear Resistant Material In non-explosive crushing methods, wear occurs when a hard rock makes an impact on the surface of the crusher, its teeth, or liners, etc. [3]. This interaction leads to the degradation of the crusher’s components, resulting in an overall decrease of the crusher’s performance [11]. Wear is defined as the degradation of a material’s surface by removing (or displacing) material by the mechanical action of a contacting substance. Abrasion wear is a type of wear where elastic deformation and surface damage occurs at hard asperities, resulting in a material loss caused by the hard asperities (Figure 2.5 A). Abrasion wear can also result from an abrasive particle sliding between two adjacent surfaces (Figure 2.5 B) [12]. Abrasion wear can be further classified as: (1) low-stress abrasion and (2) high-stress abrasion. Low-stress abrasion wear either leads to the removal of a material or forms scratches on the component’s surface [12]. Also, there is no major fracture observed (e.g. a free flow of dumped ore on the surface of the truck bed as shown in Figure 2.6 A). In high-stress abrasion wear, when large rocks are handled by the rotating crusher parts, such as teeth of roller crushers, A B Figure 2.5: Two types of abrasion wear: (A) two-body abrasion, (B) three-body abrasion [12]. 8 wear occurs [3]. In this wear condition, the material removal from the crusher’s part is by plastic deformation, as seen in Fig. 2.6 (B) [12]. Therefore, in low-stress and high-stress abrasion, the hardness of the material plays a major role to reduce the surface damage by the hard particles. Impact wear is also observed when hard rocks impact the crusher’s surface, teeth, or are stuck between the moving jaws, thereby resulting in high localized pressure [13]. This pressure can cause surface damage, because of wear and may result in plastic deformation (depending on the impact strength and size of the rocks). For impact wear applications, a material should have high hardness for wear resistance with the high impact resistance to handle the impact of rock. In conclusion, a material with high hardness (>500 HV) will experience lower abrasion wear, but a material with high toughness (>100 Joules) coupled with high hardness is more suitable for impact wear application. Material Options for Crushing Applications High Cr white iron alloys can be used for the crusher teeth or liner’s surface, since these alloys have excellent wear properties [2], [4], but were seen to fail under repetitive loading [5]. A B Figure 2.6: Picture of: (A) truck bed is abraded by the free flow of dumped ore or minerals under relatively low wearing stress, and (B) wear of roller crusher under high stress [12]. 9 Moreover, they are expensive [2], and there is a need to find an alternative material with low cost which can sustain repetitive loading. Figure 2.7 shows the toughness vs. abrasion of austenitic manganese, martensitic, white cast iron, and ceramics [14], [15], which have been often considered for abrasion resistant material applications. Although white cast iron and martensitic steels have a good wear resistance, austenitic manganese steel or Hadfield steel is considered suitable for crushing equipment [16]. Hadfield steel is work-hardened without becoming brittle under impact loading [2]. These steels are preferred for the crusher’s parts (in the primary crushing stage) [1] as martensitic and white cast irons, as shown in Figure 2.7 [14], [15], have a lower impact toughness. 2.2 Introduction to Hadfield Steels In 1882, Sir Robert Hadfield discovered a class of alloys known as Hadfield steels or austenitic/high manganese steels [4], [12], [16]. Figure 2.8 shows a jaw crusher plate made out of Hadfield steel [17]. These steels have a high toughness and ductility, good casting ability, and work hardenability [18]–[20]. The high toughness and ductility arise from the solid Figure 2.7: Toughness vs. abrasion resistance [14], [15]. 10 solution strengthening due to retained C and Mn in the austenite Fe after the heat-treatment [4]. The reason behind the good castability is the high fluidity of Hadfield steels [4], [17]. Even at low superheat, it is possible to cast complex mold shapes [17]. The work hardening ability of Hadfield steels during impact depends on various factors, such as alloy composition and strain rate [4]. There is no significant decrease in the toughness of Hadfield steels during work hardening [2], as will be discussed in detail in Section 2.6. The casting parameters, heat-treatment procedure, and chemical composition play a significant role in deciding the mechanical properties of Hadfield steels. These factors are discussed in the next sections. 2.3 Chemical Composition The chemical composition of Hadfield steels ranges from 1-1.4% carbon with approximately 10-14% Mn, with balance Fe [21]. It is a non-magnetic steel with a face-centred cubic lattice. Hadfield steel is cast based on ASTM standard A128. Table 2.1 shows the ASTM grades of Hadfield steels [4], [22]. The wear resistance of the Hadfield steels depends mainly on the Figure 2.8: Jaw crusher plate made of Hadfield steel [17]. 11 percentage of C and Mn content [6]. With an increase in wt. % of C and Mn, the formation of carbides increases, therefore, the wear resistance of the Hadfield steels increases [18]. Table 2.1: Standard ASTM grades of Hadfield steels [4], [22]. Effect of Manganese Content Mn is a strong austenite stabilizer that delays the phase transformation of austenitic Fe to ferritic Fe or martensitic Fe [4]. During the quenching of steels containing 1.1 wt. % Mn to 370 ℃, the transformation from austenite to ferrite starts in 15 sec, whereas for 13 wt. % Mn alloys it is delayed for 48 hrs [4], [18]. Figure 2.9 shows the binary phase diagram of Fe-Mn alloys. With an increase of Mn from 10 to 25 wt. %, the solidification range increases from 15 to 30 ℃ [23]. The austenitic Fe is stable up to 550 ℃ then enters the two-phase (austenitic Grade C Mn Cr Mo Ni Si (Max) P (Max) A 1.05-1.35 11.0 min - - - 1.0 0.07 B-1 0.9-1.05 11.5-14.0 - - - 1.0 0.07 B-2 1.05-1.2 11.5-14.0 - - - 1.0 0.07 B-3 1.12-1.28 11.5-14.0 - - - 1.0 0.07 B-4 1.2-1.35 11.5-14.0 - - - 1.0 0.07 C 1.05-1.35 11.5-14.0 1.5-2.5 - - 1.0 0.07 D 0.7-1.3 11.5-14.0 - - 3.0-4.0 1.0 0.07 E-1 0.7-1.3 11.5-14.0 - 0.9-1.2 - 1.0 0.07 E-2 1.05-1.45 11.5-14.0 - 1.8-2.1 - 1.0 0.07 F 1.05-1.35 6.0-8.0 0.9-1.2 - 1.0 0.07 12 Fe and ferritic Fe) region. However, the austenitic Fe phase is stable at room temperature because of the low diffusivity of atoms [24]. Figure 2.10 depicts the tensile strength, ductility, yield strength, and elongation of Hadfield steels as a function of Mn content. The mechanical properties of high Mn steels are greatly influenced by the fraction of carbides in the microstructure and retained Mn in the austenitic Fe matrix. An increase in Mn from 10 to 14 wt.% resulted in increased tensile strength and ductility due to solid solution strengthening. The tensile strength and ductility decreased with further addition of Mn up to 16 wt. %, [4], [25]. This is due to an increase in the formation of carbides [18]. The yield strength increased with the increase of Mn wt.% due to the retained Mn as a solid solution in austenitic Fe after the heat-treatment. The stabilization of austenitic Figure 2.9: Fe-Mn binary phase diagram [23]. 13 Fe and the probability of formation of carbide phases increased with an increase in Mn wt. % due to austenitic stabilizer and carbide forming properties of Mn [18]. Lencina et al. [6] investigated the effect of Mn content (12.5 and 16.6 wt.%) on the wear performance. The composition of both the alloys is shown in Table 2.2. Table 2.2: Chemical composition of alloys with 12 and 16 wt. % Mn [6]. The performance analysis of both types of Hadfield steels was evaluated in a pebble crusher. They found that an alloy with 16 wt. % Mn had an extra service time of 45 hours and a 10% increase in productivity than an alloy with 12 wt. % Mn, as shown in Table 2.3 [6]. The reason behind the high productivity was due to the high hardness of 16 wt. % Mn alloy, which resulted Alloy C Mn Cr Mo Ni Si P S Nb 12 wt. % Mn 1.46 12.50 1.46 0.10 0.20 0.53 0.030 0.001 0.12 16 wt. % Mn 1.41 16.68 1.53 0.08 0.13 0.48 0.034 0.003 0.11 Figure 2.10: Effect of Mn content on the mechanical properties and microstructure of Hadfield steels [18]. 14 in high resistance against wear. The hardness of the 16 wt. % Mn alloy was observed to be around 600-700 HV, whereas the hardness for the 12 wt. % Mn alloy was around 300 HV. Table 2.3: Performance characteristics of 12 wt.% Mn and 16 wt.% Mn steels [6]. Lencina et al. also found that the alloys with 12 wt.% Mn and 16 wt.% Mn had an average impact value of 50 J and 225 J, respectively [6]. This was attributed to solid solution strengthening from the retained Mn in the austenitic matrix in the alloy with 16 wt. % Mn [4]. Zavadil et al. studied the effect of grain boundary coverage of carbides on the toughness and found that the toughness decreased exponentially with an increase in grain boundary coverage by the carbides [26]. This may be one of the reasons behind the better performance of 16 wt. % Mn alloy (due to the discontinuous carbides in the microstructure seen in Figure 2.11 (B)). On the other hand, continuous carbide along grain boundary in Figure 2.11 (A), may lower the toughness of 12 wt. % Mn alloy. Alloy Weight Loss (%) Service Time (hrs) Productivity (t) Wear (g/t) 12 wt. % Mn 32.7 390 103000 7.84 (±0.63) 16 wt. % Mn 24.1 435 115000 5.18(±0.41) A B Figure 2.11: Micrograph of alloys with: (A) 12 wt. %Mn addition, and (B) 16 wt. % Mn addition [6]. Continuous carbides Discontinuous carbides 15 Effect of Carbon Content Carbon (C) is an austenite stabilizer and plays a vital role in the abrasion resistance of the Hadfield steels due to its ability to form (Fe,Mn)3C carbides [13], [27]. The solubility limit of C in austenitic Fe is 2 wt. %. The formation of these types of carbides increases with an increase in wt. % of C, which increases the hardness of the alloy resulting in an improvement in wear resistance [18]. In general, the tensile strength and ductility increase with the increase in C up to 1.2 wt.%, as shown in Figure 2.12 [18] because of solid solution strengthening of retained C in austenitic Fe after the heat-treatment. Above 1.2 wt. % C, carbides form because C cannot be retained in a solid solution in austenitic Fe after the heat-treatment [28]. Thus, coarse carbides form, which results in a decrease of the tensile strength and elongation, as seen in Figure 2.12. Figure 2.12: Effect of varying carbon (wt. %) on the mechanical properties and microstructure of Hadfield steels [18]. . 16 2.3.2.1 Effect of carbon on the grain size of Hadfield steel At 6.4 atomic % C (1.45 wt. % C), the vertical line in Figure 2.13 [29], shows the temperature difference of 225 ℃ between the liquidus and solidus line. This shows how the variation in carbon content affected the solidification range of the alloy which is directly linked to its properties. Figure 2.13: Phase diagram of Fe-C-Mn [29]. Gürol et al. [30] investigated the effect of varying C on the grain size of Hadfield steels containing 17 wt.% Mn. They found that with an increase in carbon content from 0.7 to 1.3 wt. % in 17 wt. % Mn steel, the average grain size increased from 281 to 581 μm. Table 2.4 shows the solidus and liquidus temperature decreased with an average of 45 ℃ and 100 ℃, 17 respectively. With an increase in C from 0.7 to 1.3 wt.%, the solidification range of the alloys increased by approximately 1.5 times. This led to an increase in the grain size of the alloys with an increase in the freezing range as shown in Table 2.4 [30]. Table 2.4: Dependence of alloy grain size on the solidus and liquidus temperatures [30]. Alloy Solidus temperature (℃) Liquidus temperature (℃) Freezing range (℃) Average grain size (μm) 0.7C-17Mn 1305 1405 100 281 ± 18 1.1C-17Mn 1255 1380 125 338 ± 26 1.3C-17Mn 1205 1360 155 589 ± 38 Effect of Other Alloying Elements Several alloying elements such as Cr, Si, Mo, P, and S are also added to Hadfield steels. The contribution of each element is discussed below. 2.3.3.1 Chromium Chromium (Cr) is added to Hadfield steels to increase its yield strength and resistance against corrosion and abrasion [4], due to the strong ability of Cr to form chromium carbides as compared to Fe and Mn [31]. The probability of formation of (Fe, Cr, Mn)23C6 type complex carbides and (Fe,Mn)3C increases with an increase in wt.% of Cr [18] because of stabilizing effect of Cr on Fe based carbide. With an increase of Cr up to 2.5 wt. % in the Hadfield steel, the yield strength increased due to solid solution strengthening [4]. Tensile strength and elongation increased with an increase in Cr up to 2 and 1.7 wt.%, respectively. Then these values (tensile strength and elongation) decreased with further addition of Cr, as shown in Figure 2.14 [18]. This can be due to an increase in the volume fraction of carbides in the microstructure with an increase in wt. % of Cr, which is detrimental to the tensile strength and elongation [31]. 18 Chen et al. [28] studied the effect of Cr and N on the Hadfield steels containing approximately 11.6 wt. % Mn. The composition of steels containing around 11.6 wt. % Mn with and without Cr are shown in Table 2.5. Table 2.5: Chemical composition of 12 wt. % Mn steels with and without Cr [28]. It was observed that there was no significant difference in the grain size and hardness (of steel containing 2.14 wt. % Cr) when the austenitizing temperature was varied from 900 to 1200 ℃, as shown in Figure 2.15. On the other hand, in an alloy with 11.6 % Mn, the grain size increased with an increase in the austenitizing temperature. The grain size of 11.6 % Mn 2.14 Steels C Mn Cr N Si Al Cu P S 11.6 % Mn 2.14 %Cr 1.10 11.6 2.14 0.052 0.42 0.024 0.014 <=0.001 0.005 11.6 % Mn 1.14 11.7 0.07 0.016 0.42 0.027 0.009 <=0.001 0.005 Figure 2.14: Effect of varying chromium (wt.%) on the mechanical properties and microstructure of Hadfield steels [18]. 19 %Cr steel was observed to be finer than 11.6 % Mn steel because of the uniform distribution of Cr and N in the microstructure and this could have impeded the grain boundary movement for grain coarsening [28]. The hardness of steel containing 2.14 wt. % Cr was observed to be 252 HV after solution heat-treatment at 1100 ℃, whereas the hardness of steel without Cr reached only 231 HV. This was due to an increase in the hardness with an increase in C wt. %, and the formation of carbides due to Cr in Hadfield steels [4]. 2.3.3.2 Silicon Si is a carbide forming element and is added to Hadfield steels up to 0.8 wt. % concentration for better fluidity during casting and to improve the hardness and wear resistance [18]. The addition of Si up to 1 wt.% is considered safe, because Si does not form needle shape SiC carbides, which may affect the mechanical properties of the Hadfield steels [18]. Addition of Si from 1 to 2 wt. % is preferred for increasing the yield strength of the Hadfield steels [4]. Above 2 wt.% Si, the strength and ductility of the Hadfield steels significantly decreases due to the formation of the needle shape carbides [4]. Figure 2.15: Variation in grain size (µm) and hardness (HV) for 11.6 wt. % Mn alloy with and without Cr [28]. 11.6 % Mn 2.14 %Cr 11.6 % Mn 2.14 %Cr 11.6 % Mn 11.6 % Mn 20 2.3.3.3 Molybdenum Molybdenum (Mo) is an austenite stabilizer, forms MoC carbides, and promotes the dispersion of discontinuous carbides in the austenitic Fe grains during a heat-treatment [17]. The microstructure with dispersed carbides improves the wear resistance and toughness of the Hadfield steels [4]. Therefore, the addition of Mo is beneficial in the casting of heavy sections (thickness >10 inches) of the crusher’s part because of the dispersion of carbides. The addition of Mo above 1 wt.% may result in decreasing the mechanical properties due to the incipient melting phenomenon [17]. It is a melting phenomenon that occurs due to the presence of low-melting constituents inside and along the grains during the heat treatment. 2.3.3.4 Nickel Ni is a strong austenitic stabilizer that also delays the transformation of the austenitic Fe to ferritic Fe or martensitic Fe [4], [17]. It also suppresses the formation of carbides even at a low quenching rate in heat-treatment [17]. The addition of Ni increases the ductility and lowers the abrasion and yield strength of the Hadfield steel [4]. Ni is usually added to the welding grade (low C grade) of High Mn steels to suppress the formation of carbides [17]. 2.3.3.5 Phosphorus and Sulphur Phosphorus forms a phosphoid eutectic film at the grain boundaries, which results in harming the mechanical properties of the steel [17]. Sulphur generally does not have any effect on the mechanical properties because of the scavenging effect shown by Mn. In this effect, Mn leads to an elimination of S, by forming harmless and rounded sulphide inclusion in the microstructure [4]. However, in practice, it is recommended to keep P and S concentrations as low as possible due to their ability to form the phosphoid eutectic film and sulphide inclusions, respectively. 21 2.4 Casting Parameters The casting process for Hadfield steels is similar to other carbon steels [32] and molten material for the casting can be prepared by using an electric arc [4] or induction furnace [18]. It is possible to fill the complex mold shapes even with a low superheat because of the high fluidity of high Mn steels [17]. The pouring temperature should be maintained below 1500 ˚C to reduce the slag formation that might take place between Mn and refractory liners [25]. The selection of mold material is one of the essential parts of the casting process. The acidic nature of refractory materials for molds should be avoided, because Hadfield steels strongly react with these materials and cause severe burning of silica sand [32]. Olivine sand, which is basic and free of silica, is preferred for mold preparation to prevent the reaction between the metal and the mold material [13]. Effect of Pouring Temperature Balogun et. al [25] studied the effect of three different pouring temperatures (1550 ℃, 1450 ℃, and 1380 ℃) on the microstructure of Fe-12Mn-1.23C-2.40Cr alloy. For all the three temperatures considered, as-cast materials were heat-treated at 1050 ℃ for 4 hours followed by water quenching. The authors found that higher segregation of Mn and Cr carbides was present in the microstructures of alloys poured at 1550 ℃ (Figure 2.16). The amount of segregation decreased in the microstructure of alloys poured at 1450 ℃ (Figure 2.17), whereas, a uniform distribution of carbides in the austenitic Fe matrix, with fewer carbides at grain boundaries, was seen in the microstructure of alloys poured at 1380˚C (Figure 2.18). The formation of carbides varied with the changing pouring temperature, because of an increase in 22 the solidification time, which resulted in cooling the alloy to solid from higher superheat temperature [25], as shown in the Fe-Mn phase diagram in Figure 2.9. Figure 2.18: Microstructure of 12% Mn steel poured at 1380 ℃: (A) edge, (B) middle, and (C) mounting [25]. Figure 2.17 Microstructure of 12% Mn steel poured at 1450℃: (A) edge, (B) middle, and (C) mounting [25]. Figure 2.16: Microstructure of 12% Mn steel poured at 1550℃: (A) edge, (B) middle, and (C) mounting [25]. 23 Balogun et al. [25] also carried out EDS analysis on the same alloys and quantified the degree of segregation of Mn, Cr, and Si at different pouring temperatures (1380-1600 ℃), as shown in Figure 2.19. It was concluded that there was no Si segregation due to varying pouring temperature between 1380 ℃ to 1600 ℃, while Cr and Mn segregation increased as the pouring temperature increased [25]. Unfortunately, the authors did not explain the driving force behind the segregation kinetics with the variation in pouring temperature. 2.5 Heat-treatment of Hadfield Steels Due to the high Mn content in Hadfield steels, as-cast materials readily form carbides on the grain boundaries [4]. The resulting microstructure tends to be brittle [33]. Therefore, heat-treatment becomes an essential part to dissolve the carbides [34]. This is carried out by solution heat-treatment above the Acm temperature (the temperature that corresponds to the boundary between the cementite-austenite and the austenite phase) followed by water quenching [13], [35]. The variation in heat-treatment can be optimized for obtaining the desired properties such Figure 2.19: Effect of pouring temperature on the degree of segregation of Mn, Cr and Si [25]. 24 as yield strength, tensile strength, impact strength, and abrasion resistance. Figure 2.20 shows the optical micrographs of ASTM A128 Grade B-3 before and after heat-treatment at 1120 ℃ followed by water quenching [4]. It can be seen that the carbides which led to the embrittlement have been successfully dissolved post-heat-treatment. Tecza et al. studied the effect of heat-treatment temperature (1100 to 1200 ℃) on the GX170MN16H alloy (Fe-1.7C- 16Mn-1.4Si- 1.4Cr) with varying the holding time (up to 240 minutes) in a furnace followed by water quenching [36]. They found that increasing the temperature as well as holding time did not help to achieve the pure austenite structure as shown in Figure 2.21, although a minor increase in the hardness value was observed. This may be due to the high C wt. % which increased the probability of formation of stable carbides [18]. A general heat-treatment of Hadfield steels involves slow heating to solution heat-treatment temperature with holding time for 1 to 2 hours per 25 mm thickness followed by water quenching [4]. The quenching rate plays a significant role in achieving the desired A B Figure 2.20: Optical micrographs of ASTM A128 grade B-3 steel at 500X magnification : (A) before heat-treatment, and (B) after heat-treatment at 1120 ℃ [4]. 25 microstructure. It is hard to increase the cooling rate during the quenching due to the low thermal conductivity of the alloy [37]. In addition, the cooling rate is greatly influenced by the thickness of the casting. When the cast thickness was more than 100 mm, strong segregation of chromium and manganese carbides was observed due to the low thermal conductivity of Hadfield steels [36]. 2.6 Work-hardening of Hadfield Steels Many debatable theories have been proposed by various authors on the work hardenability of Hadfield steels. Tasker et al. [13] have reported that the hardness of the Hadfield steel increased from ~200 to ~650 HV due to work-hardening on the worn surface. Some researchers have proposed that during loading, a phase transformation in the alloy occurs from austenitic Fe to ferritic Fe or martensitic Fe [13]. On the contrary to this statement, grade A to E of Hadfield steels (shown in Table 2.1) have never shown any phase transformation under loading conditions. During the plastic deformation, Hadfield steels with Mn greater than 10% were stable even below 0℃ temperatures [38]. In ASTM grade F of Hadfield steels, in which A B Figure 2.21: Microstructure of cast Hadfield steel after solution treatment followed by water quenching with water maintained at: (A) 1200 ℃ with 80 minutes of hold time, and (B) 1150 ℃ with 240 minutes of hold time [36]. 26 Mn content is 6-8 wt. %, this transformation from austenitic Fe to ferritic Fe or martensitic Fe during loading has been observed [13]. The mechanisms resulting in the strain hardening involve the twining (single or multiple) [17], the interaction between the interstitial carbon atom and dislocation [39], [40], and the interactions between dislocations and twining mechanism [41], [42]. Raghavan et al. [39] connected the rate of work hardening to mechanical twinning. They found that the twin boundaries hindered the movement of dislocations and resulted in the work hardening of Hadfield steels. On the contrary to this study, Lambakahar et al. [43] observed that the work hardening in an alloy with coarse grain size was less influenced by twinning. This is because of the low density of dislocations in steels with coarse grains. This resulted in less hindrance for twin dislocation, which further led to a smaller work hardening coefficient. Karanam et al. [41], [42] studied the single crystal material and observed that the activation of multiple twinning systems resulted in a higher strain hardening coefficient than a single twin system. The dislocation interaction with C and Mn atoms played a major role in the work hardening mechanism of Hadfield steels. Yan et al. [44] studied the work hardening rate using the shot peening method and observed the increase in the hardness of the surface as the shot peening time increased, as seen in Figure 2.22. The hardness increased from 250 to 750 HV after 120 min of shot peening. This was due to the development of high-density dislocations and twin formation. 27 2.7 Chapter Summary In the current literature, most of the work was done on studying the effect of casting and heat-treatment parameters on the microstructure, hardness, and impact toughness of Hadfield steels. A small number of papers dealt with the work hardening mechanism of Hadfield steels. Many researchers have worked on studying the effect of varying C, Cr, Si, Mo, Ni, and P in the Hadfield steels. Studies on the microstructure, phases, wear performance and toughness of high Mn steel were limited to 20 wt. % Mn. Thus, the current research aims to focus on the effect of varying the Mn content up to approximately 23 wt. %. The microstructure analysis, phases phase analysis, toughness (low and high strain), hardness, and wear resistance were studied to fill the gap in the current literature. Figure 2.22: Variation in hardness (HV) vs. shot peening duration [44]. 28 Chapter 3: Experimental Procedure This chapter provides an overview of the experimental procedures including the casting and heat-treatment processes, microstructure evaluation, phase analysis, and mechanical testing. 3.1 Casting and Heat-treatment Process Three different compositions of Hadfield steels with a varying Manganese content were cast at Spokane Industries, Spokane Valley, Washington, United States. A typical high Mn steels has 10- 15 wt. % Mn as per ASTM A128 [4], [30]. Therefore, the two compositions with a step size of 5 wt. % from the standard Mn content were prepared in the present work. Table 3.1 shows the composition of the as-cast material. Based on the concentration of Mn content, alloy designations were assigned as follows:12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn. Castings were made using sand molds. Table 3.1: Chemical composition of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn. The degree of segregation of carbide forming elements such as Si, Cr, and Mn increases with an increase in tapping temperature (1380-1600℃) of High Mn steels as discussed in Section 2.4. The tapping temperature for steels is usually 25-30℃ above the melting point to compensate for the heat loss during the pouring process. Therefore, the tapping temperature was set at 1415℃. Figures 3.1 (A) and (B) show the front and top views of the as-cast alloys, Alloy C Mn Cr Si P Ni Mo Cu Fe 12 wt. % Mn 1.05 12.14 2.00 0.74 0.02 0.13 0.03 0.17 83.68 18 wt. % Mn 1.13 17.70 1.99 0.40 0.01 0.06 0.02 0.14 77.25 23 wt. % Mn 1.15 22.78 2.06 0.71 0.01 0.10 0.02 0.13 72.97 29 respectively. After casting, the alloys were heated to 1065 ℃ as shown in Figure 3.2 and held isothermally for 4 hours to dissolve the carbides in austenitic Fe. Afterward, the alloys were quenched in water at 21 ℃. Figure 3.2: Heat-treatment chart of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn. B A Figure 3.1: Pictures of casting block: (A) front-view, and (B) top-view (all dimensions in cm). 6.9 10.1 22.8 0200400600800100012000 2 4 6 8Temperature (℃)Time (hrs)30 3.2 Microstructure Analysis The microstructure of the alloys was analyzed using optical microscopy and scanning electron microscopy techniques. Sample Preparation Three cylindrical samples of diameter 3.81 cm and length 2 cm were machined using the electrical discharge machining from 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn casting blocks at JAS Industrial, Calgary, Canada. The cylindrical samples were polished on a manual polisher using 120, 240, 320, 400, and 600 grit SiC papers to achieve a flat surface. Figure 3.3 shows the optical sample. Following this, the samples were polished on a polishing pad with diamond pastes of 9, 6, 3, and 1 µm. The samples were ultrasonically cleaned for 5 minutes after each step using an ultrasonic bath. The samples were cleaned with alcohol to clean the surface and then etched by immersing in 2% nital solution (2 ml nitric acid mixed with 98 ml ethanol) for 90 seconds followed by an alcohol bath. Figure 3.3: Picture of cylindrical microscopy sample. Diameter=3.81 cm 31 Optical Microscopy Zeiss optical stereoscopic microscope (Axiovert Alm, model 37081) was used for the optical microscopy of alloys. 3.2.2.1 Grain Size Measurement The average grain size of the alloys was calculated using the Heyn linear intercept method as per ASTM E112-10 [45]. ImageJ software was used to measure the intercept lines. The optical microscope was used to capture micrographs representing the general microstructure of the alloys. Single or multiple straight lines with more than 50 intercepts in a single field were used for measuring the average grain size of the alloys. The average grain size was calculated by yielding more than 400 intercepts to attain 10% relative precision accuracy. Scanning Electron Microscopy Figure 3.4 shows the Tescan® Mira3 XMU field emission scanning electron microscope present at UBC Okanagan Filter laboratory. This setup was used for imaging in secondary electron mode. Figure 3.4: Field emission scanning electron microscopy setup. 32 3.2.3.1 Elemental Analysis The elemental composition and phase analysis of alloys was performed using Elemental X-ray Energy Dispersive Spectroscopy (X-EDS) and XRD diffraction respectively. The elemental analysis for each material was performed using Oxford® AZtec X-MAX X-Ray Energy Dispersive Spectroscopy (X-EDS) connected to the SEM microscope. The samples prepared for optical microscopy were used for elemental analysis as well. 3.3 Phase Analysis X-Ray Diffraction XRD analysis was carried out at the X-Ray Crystallography Facility at the Department of Chemistry, University of British Columbia, Vancouver. A Bruker D8 Advance X-ray Diffractometer was used for all measurements with a summary of the experiment parameters listed in Table 3.2. The phases were identified by analyzing the qualitative results using the X'Pert High Score software. Table 3.2: XRD test parameters. Parameter Description Sample Size 2.54 cm x 2.54 cm square with thickness of 0.3175 cm Radiation Copper Kα1 & Kα2 Filter Nickel (filter out CuKβ) Detector LynxEye silicon strip Generator 40 kV, 40mA Slits 1mm divergent, 0.8 cm anti-scatter, 2.5˚soller Step size 0.03˚ Step time 336 sec Discriminator Settings LL= 0.19 and WW=0.06 33 3.4 Mechanical Property Evaluation The mechanical properties of the investigated materials were analyzed using Charpy-impact, tensile, pseudo-wear, and microhardness testing. The procedure used for the sample preparation and testing for each of these tests is discussed below. Charpy-Impact Testing 3.4.1.1 Sample Preparation The samples for the Charpy-impact were prepared as per ASTM standard E23-18b [46]. The V-notch samples were cut to the standard size as shown in Figure 3.5 using electrical discharge machining. The three samples were prepared from each alloy composition. Figure 3.5: V-notch Charpy-impact sample size [46]. 34 Charpy-impact testing was carried on the Charpy-impact tester shown in Figure 3.6 available at the School of Engineering, UBC Okanagan. The value of Charpy-impact toughness reported in this thesis represents an average of three measurements. Tensile Testing 3.4.2.1 Sample Preparation The samples for the tensile testing were prepared as per ASTM standard E8/E8M-13 [47]. The sub-size sample geometry, as shown in Figure 3.7 (A), was used due to material constraints. The samples were cut using the electrical discharge machining and three samples per composition were prepared for testing. Figure 3.7 (B) shows the prepared tensile sample. Figure 3.6: Charpy-impact tester. 35 The INSTRON Tensile test machine (Model 5969, 50 KN) at Composites Research Network, UBC Okanagan was used for testing. The tensile testing was conducted at a strain rate of 1.38 mm/min. The average values for 0.2% offset yield strength, ultimate tensile strength, and % elongation were calculated based on three samples for alloys with 12 wt. % Mn and 18 wt. % Mn, and two samples for an alloy with 23 wt. % Mn. Pseudo-Wear Testing Cylindrical samples of diameter 3.2 cm and 1.4 cm length were prepared using electrical discharge machining. Two samples for each composition were prepared for testing. A pseudo-wear test was performed on the samples using auto-polisher equipment. In this test, the samples were abraded against a 120 grit SiC paper with fixed force and time. The weight loss of the samples was measured to provide a relative measure of the wear resistance of the different materials. A Struers Tegramin-30 autopolisher, as shown in Figure 3.8 (A), was used for the pseudo-wear testing. The flow rate of water for the test was set manually to 19.23 ml/sec for the testing. The total piston force was set to 330 Newton (55N/piston) and the speed of the disc and head were Figure 3.7: Picture of: (A) sub-size tensile sample dimensions [47], and (B) prepared tensile sample. A B 36 set to 150 RPM clockwise. The samples were placed from P1 to P6 position in the sample holder, as seen in Figure 3.9. Further, the samples placed on P1 to P6 position were aligned with relevant positions marked from P1 to P6 on the Auto-polisher head, as seen in Figure 3.8 (B). The samples were rotated one step clockwise from their position after each cycle run. Table 3.3 shows the stages of testing with different testing and cycle run times. Each sample’s weight loss was measured to calculate the weight loss % of each composition after each cycle run. Disc with 120 Grit Paper Samples holder Head with six pistons Water Outlet A B Figure 3.8: Picture of: (A) Struers Tegramin-30 auto-polisher, and (B) samples mounted on auto-polisher. Figure 3.9: Sample holder with position P1 to P6. 37 Table 3.3: Stages of pseudo-wear test. Microhardness Testing The samples used for the pseudo-wear testing were also used for measuring the average microhardness. The hardness was measured using the Micro Vickers Hardness Tester model SXHV-1000TA (shown in Figure 3.10) with a load and dwell time of 9.807 Newton and 5 seconds, respectively. The test was performed as per the ASTM E91-17 standard [48]. A minimum of 20 measurements was used to calculate the average hardness of each material. The hardness value of the samples after each stage of the pseudo-wear test was also measured. Stage Testing time (min) Cycle run time (min) 1 12 1 2 60 5 3 60 5 4 60 5 Figure 3.10: Microhardness tester. 38 Chapter 4: Results and Discussion This chapter presents the results obtained from the Fe-Mn-C system, microstructure, elemental and phase, and mechanical analysis of this thesis work. 4.1 Fe-Mn-C System Analysis The Thermo-Calc software with Fe-demo database was used to create the Fe-Mn-C liquidus projection and phase diagram to study the effect of Mn in high Mn steels. Liquidus Projection Figure 4.1 shows the liquidus projection of Fe-Mn-C ternary system generated using Thermo-calc software. During the casting process, the alloy is heated above melting point to compensate for the loss of heat and to fill complex mold shapes. The difference between the melting temperature and pouring temperature is defined as superheat. For high Mn steel, it is possible to fill the complex mold shapes even with a low superheat because of its high fluidity [17]. Balogun et. al [25] studied the effect of pouring temperature on the degree of segregation of carbide forming elements (such as Si, Cr, and Mn) in alloy with 12 wt. % Mn. They found that with an increase in temperature from 1380 to 1600℃ of Fe-12Mn-1.23C-2.40Cr alloy, the degree of segregation of Mn and Cr increased exponentially due to an increase in superheat. Therefore, the tapping temperature was set at 1415℃ for the casting of the present alloys to reduce the degree of segregation. The liquidus temperature of Fe-12.14Mn-1.05C, Fe-17.7Mn-1.13C, and Fe-22.78Mn-1.15C alloys are marked in Figure 4.1. The liquidus temperature of all the alloys were calculated using the linear interpolation. Table 4.1 shows the summary of the calculated liquidus temperature (°K and ℃) and superheat for all the alloys. The liquidus temperature decreased from 1412.5℃ to 1370℃ with an increase of Mn from 12.14 to 22.78 39 wt. %. On the other hand, the calculated superheat increased by 18 times with an addition of Mn from 12.14 to 22.78 wt. %. The increase in super heat could result in the degree of segregation of carbide forming elements which further results in uneven wear properties of high Mn steels. Therefore, the superheat should be adjusted by lowering the tapping temperature with respect to Mn wt. % percentage in high Mn steel to reduce the segregation of carbide forming elements. Fe-12.14Mn-1.05C (1685.65°K) Fe-17.7Mn-1.13C (1661.15°K) Fe-22.78Mn-1.15C (1643.15°K) 1643.15° 1633.15°K Figure 4.1: Liquidus projection of Fe-Mn-C system (all temperatures in Kelvin). 40 Table 4.1: Liquidus temperature and superheat of Fe-12.14Mn-1.05C, Fe-17.7Mn-1.13C, and Fe-22.78Mn-1.15C alloys. Phase Analysis Figures 4.2-4.3, 4.4-4.5, and 4.6-4.7 show the ternary Fe-Mn-C equilibrium phase diagrams at 900℃, 750℃, and 650℃, respectively. The volume percentage of cementite was calculated using the volume fraction of phase tool in Thermo-Calc. The volume fraction of cementite can be seen for the marked alloys Fe-12.14Mn-1.05C, Fe-17.7Mn-1.13C, and Fe-22.78Mn-1.15C in Figures 4-3, 4-5, and 4-7. Table 4.2 shows the summary of the calculated volume percentage of cementite of all the compositions at 900℃, 750℃, and 650℃, respectively. As shown in Table 4.2, the volume percentage of cementite increased for all the compositions with a decrease in isothermal section temperature from 900℃ to 650℃. Also, the volume percentage of cementite increased with an increase in Mn in the alloys at 650℃ and 750℃. Similarly, Gurol et.al [30] observed an increase in volume fraction of carbides from 7.5 to 18.6 % with varying Mn from 13 to 20 wt. % Mn in as-cast high Mn steels, respectively. The ɣ-(Fe, Mn) phase (austenitic Fe with Mn in solid solution) appeared in the isothermal section of Fe-Mn-C at 900℃, in Figure 4.3, for all the marked compositions. With a decrease in temperature from 900℃ to 750℃ and 650℃ for the Fe-Mn-C system, the alloy's composition had ɣ-(Fe, Mn) and cementite phases, as seen in Figures 4.5 and 4.7, respectively. Composition Liquidus Temperature (°K) Liquidus Temperature (℃) Superheat = 1415℃ - Liquidus Temperature (℃) Fe-12.14Mn-1.05C 1685.65 1412.5℃ 2.5 Fe-17.7Mn-1.13C 1661.15 1388℃ 27 Fe-22.78Mn-1.15C 1643.15 1370℃ 46 41 In plain carbon steels, the ɣ-Fe begins to decompose at 723℃ (eutectoid temperature) to form α-Fe. As seen in the Fe-Mn-C phase diagram at 650℃ in Figure 4.7, there was no α-Fe for the alloy with Mn above 10 wt. % with C in the range 0-1.7 wt. %. This is due to the austenitic stabilizing properties of Mn which lowers the phase transformation temperature and suppresses the formation of α-Fe by delaying the process [4], [18]. Table 4.2: Volume percentage of cementite in Fe-12.14Mn-1.05C, Fe-17.7Mn-1.13C, and Fe-22.78Mn-1.15C at 650℃, 750℃, and 900℃. Compositions Volume percentage of Cementite at 900 ℃ Volume percentage of Cementite at 750 ℃ Volume percentage of Cementite at 650 ℃ Fe-12.14Mn-1.05C 0 5.34 9.5 Fe-17.7Mn-1.13C 0 6.73 11.03 Fe-22.78Mn-1.15C 0 7.03 11.54 Figure 4.2: Ternary phase diagram of Fe-Mn-C system at 900 °C. 42 * ⋄ º Fe-17.7Mn-1.13C (0 volume % Cementite) Fe-22.78Mn-1.15C (0 volume % Cementite) Fe-12.14Mn-1.05C (0 volume % Cementite) 0 % Cementite) 900 ℃ Figure 4.3: Ternary phase diagram of Fe-Mn-C system at 900 °C (Mn: 0-50 mass %). Figure 4.4: Ternary phase diagram of Fe-Mn-C system at 750 °C. 43 * ⋄ º Fe-12.14Mn-1.05C (5.34 volume % Cementite) Fe-17.7Mn-1.13C (6.73 volume % Cementite) Fe-22.78Mn-1.15C (7.03 volume % Cementite) 750 ℃ Figure 4.5: Ternary phase diagram of Fe-Mn-C system at 750 °C (Mn: 0-50 mass %). Figure 4.6: Ternary phase diagram of Fe-Mn-C system at 650 °C. 44 Figures 4.8 and 4.9 show the Fe-Mn-C ternary phase diagram at 500℃. In Fe-Mn-C isothermal section at 500℃, the α-Fe, ɣ-(Fe, Mn), and cementite phases appeared for Fe-12.14%Mn-1.05%C and Fe-17.7%Mn-1.13%C alloys, whereas the ɣ-(Fe, Mn) and cementite formed for Fe-22.78%Mn-1.15%C alloy. On the other hand, in Fe-Mn-C isothermal section at 250℃ in Figures 4.10 and 4.11, the α-Fe phase appeared in all three compositions. High Mn steel is also known as austenitic Mn steel due to its austenitic Fe (ɣ-Fe) microstructure and non-magnetic property [4], [18]. The addition of Mn to the steels helps to delay the transformation of ɣ-Fe to α-Fe due to its austenitic stabilization property [21]. As per ASTM standard A128/A128M, the standard grade A to E of high Mn steels (Mn: 11-14.0 wt. %) have austenitic Fe microstructure [22]. Gurol et. al [30] investigated the effect of Mn variation from Figure 4.7: Ternary phase diagram of Fe-Mn-C system at 650 °C (Mn: 0-50 mass %). 45 13 to 20 wt. % and found the ɣ-Fe and cementite phases in the XRD analysis of the alloys in the as-cast condition. Therefore, the presence of α-Fe in the Fe-Mn-C ternary phase diagrams at 250℃ for Mn > 10% could be due to the equilibrium conditions in which Thermo-Calc software allowed sufficient time to homogenize the composition. In practice, the better cooling rate does not allow the homogenization of composition simultaneously during the process. The industry-relevant cooling rates and suppressed diffusion of elements in solid phases delay the gamma-alpha phase transformation, leading to the predominantly-gamma (or fully gamma) phase structure at the low temperatures. Figure 4.8: Ternary phase diagram of Fe-Mn-C system at 500 °C. 46 Figure 4.9: Ternary phase diagram of Fe-Mn-C system at 500 °C (Mn: 0-50 mass %). Figure 4.10: Ternary phase diagram of Fe-Mn-C system at 250 °C. 47 4.2 Microstructure Analysis Optical Microscopy Figure 4.12 shows the general microstructure of an alloy with 12 wt. % Mn. The microstructure of this alloy primarily consisted of austenitic grains with few intermittent carbides at the grain boundaries, as seen in Figures 4.12 (C) and (D). The continuous carbides form readily around the grain boundaries during the solidification in Hadfield steels and make the alloys brittle [4]. Therefore, heat-treatment becomes essential to dissolve the carbides and to retain the C in a solid solution to retain the austenite Fe microstructure free of carbides [34]. Figure 4.11: Ternary phase diagram of Fe-Mn-C system at 250 °C (Mn: 0-50 mass %). 48 Figure 4.13 shows the general microstructure of an alloy with 18 wt. % Mn. As can be seen, the austenitic grains in this alloy had nearly continuous grain boundary carbide delineations, with intermittent agglomerations of the carbides. Some of the carbides had a lamellar appearance (carbides sandwiched between austenite Fe) and clusters of carbides at the grain boundaries. The accumulation of carbides on the grain boundaries has been associated with embrittlement in the high Mn steel [49], [50]. At a few sites in Figure 4.13, the carbides covered more than 50% boundary grain. A B C D Figure 4.12: Optical micrographs of an alloy with 12 wt. % Mn at: (A) 50X, (B) 100X, (C) 500X, and (D) 500X magnification. 49 Figure 4.14 shows the microstructure of an alloy with 23 wt. % Mn. The micrographs depict the carbides were present on the grain boundaries and intergranular regions. It was observed that the carbides in this alloy had a lamellar structure. The carbides inside the austenite grains can be seen in Figures 4.14 (A), (B), (C), and (E). C A B Figure 4.13 Optical micrographs of an alloy with 18 wt. % Mn at: (A) 50X, (B) 100X, and (C) 500X magnification. 50 Additional micrographs of alloys 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn can be found in Appendix A. Carbides such as M3C and M23C6 (where M is a metallic atom) are known to form during the solidification of Hadfield steel grade containing Cr, due to the solute rejection of C from the austenitic Fe [18]. It is difficult to retain all the C atoms in the solid solution after the heat-treatment due to the low cooling rate during quenching because of low thermal conductivity [4]. Gürol et al. observed that the increase in wt. % of C and Mn, the evolution of carbides with a two-phase lamellar structure in the alloys [30]. The intermittent carbides on grain boundaries appeared in all the alloys, whereas intergranular carbides and continuous C F A B E Figure 4.14: Optical micrographs of an alloy with 23 wt. % Mn at: (A) 50X, (B) 100X, (C) 500X, (D) 500X , (E) 500X, and (F) 500X magnification. D 51 carbide clusters appeared in the alloys with 18 wt. % Mn and 22 wt. % Mn. This may be due to a higher percentage of C and Mn in alloys with 18 wt. % Mn and 23 wt. % Mn than in an alloy with 12 wt. % Mn as shown in Table 4.3, since increasing the Mn and C in Hadfield steels promotes the formation of (Fe,Mn)3C carbides [18]. Table 4.3: Wt.% of C and Mn in alloy with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn addition. Alloy C wt.% Mn wt.% 12 wt. % Mn 1.05 12.14 18 wt. % Mn 1.13 17.70 23 wt. % Mn 1.15 22.78 4.2.1.1 Grain Size The average grain size was calculated using more than 400 intercepts to attain 10% relative accuracy. Single or multiple straight lines with more than 50 intercepts in a single field were used for alloys with 12% Mn and 18 wt. % Mn. The average grain size measurement of an alloy with 23 wt. % Mn was done using a collage of micrographs to yield at least 45 intercepts due to the large grain size and limited magnification options, a collage of micrographs was used. The average grain size in the alloys increased with a simultaneous increase in the Mn wt.%, as seen in Table 4.4. The pouring temperature and solidification range of Hadfield steels is one of the major factors influencing the average grain size. The pouring temperature for all the alloys was 1415 ℃. The solidification range of Hadfield steels is primarily influenced by the wt.% of C. Gürol et al. [30] found that with an increase in the C wt.% from 0.7 to 1.3 in Hadfield steels containing 17 wt. % Mn, the freezing range of the alloys increased, which resulted in grain growth as discussed in detail in section 2.3.2.1. The higher C wt. % in alloys 52 with 18 wt. % Mn and 23 wt. % Mn (shown in Table 4.4) could be one of the reasons for the grain coarsening in comparison to an alloy with 12 wt. % Mn. Mn is an austenite stabilizer that delays the transformation of γ-Fe to α-Fe [4]. With the addition of Mn content above 12 wt. %, it lowered the austenitic transformation temperature to retain austenitic Fe, as discussed in detail in section 2.3.1. . Gürol et al. observed that with an increase of Mn from 12 to 20 wt. % Mn in the Hadfield steels, it decreased both the solidus and liquidus temperature by 5 ℃ [30]. They also observed an increase in grain size from 215 to 619 µm with varying Mn from 12 to 20 wt. % in high Mn steels containing 1 wt. % C. In conclusion, there was no change in the solidification range due to Mn addition, but grain size increased. This could be one of the reasons for the grain growth of 23 wt. % Mn alloy in comparison to alloys with 12 and 18 wt. % Mn [30]. Table 4.4 Average Grain Size in alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn addition. Alloy C wt.% Mn wt.% Average Grain Size (µm) 12 wt. % Mn 1.05 12.14 116.6 18 wt. % Mn 1.13 17.70 175.8 23 wt. % Mn 1.15 22.78 445.6 4.2.1.2 Phase Area Fraction The area fraction of phases (austenitic Fe and carbides) was performed using the Fiji ImageJ software. The color threshold tool was used to calculate the area fraction. Table 4.5 shows the average phase area fraction of the alloys. Figures 4.15, 4.16, and 4.17 show the color threshold microstructure of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. Mn %, respectively. The D 53 average area fraction of carbides and austenite was calculated after performing the color threshold at eight different fields at 50X magnification for each alloy. Table 4.5: Phase area fraction in alloys with 12 wt. %, 18 wt. %, and 23 wt. % Mn addition. Alloy C wt. % Mn wt. % Carbide’s area fraction % 12 wt. % Mn 1.05 12.14 2.2 ± 0.8 18 wt. % Mn 1.13 17.70 5.0 ± 0.7 23 wt. % Mn 1.15 22.78 4.2 ± 0.9 B A Figure 4.16: Optical micrographs of an alloy with 18 wt. % Mn: (A) without color threshold, and (B) with a color threshold. Figure 4.15: Optical micrographs of an alloy 12 wt. % Mn: (A) without color threshold, and (B) with a color threshold. A B 54 Figure 4.17: Optical micrographs of an alloy with 23 wt. % Mn: (A) without color threshold, and (B) with a color threshold. One of the possible reasons for the high fraction of carbides in steels with high Mn wt. % could be an increase in superheat with Mn wt. %, as shown in Table 4.1. The degree of segregation of carbide forming elements increases with an increase in superheat, which results in carbide formation [25]. After the solution heat-treatment (SHT) at 1065 ℃ followed by water quenching, elements such as Mn, C, and Cr were retained inside the austenite grain as a solid solution, resulting in the lowest area fraction of carbides in an alloy with 12 wt. % Mn. It is difficult to retain all the C atoms in the solid solution after the heat-treatment due to the low cooling rate during quenching because of low thermal conductivity [4]. Gurol et al. [30] observed an increase in carbides fraction with an increase in C (0.7-1.3 wt. %) in an alloy with 17 wt. % Mn. Therefore, the area fraction of carbides (%) in the microstructure was the highest in alloys with 18 wt. % Mn and 23 wt. % Mn. Cr is a strong carbide forming element which forms complex M23C6 carbides (M: metallic atom) [18] and also increases the thermal stability of the carbides [37]. Therefore, a higher SHT temperature (1060-1140 ℃) was required to ensure the dissolution of carbides during heat-treatment [26]. A B 55 Scanning Electron Microscopy - General Microstructure Figures 4.18-4.20 show the scanning electron micrographs of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn. The carbides accumulated at the grain boundaries of an alloy with 12 wt. % Mn can be seen in Figure 4.18 (A) and (C). Figure 4.18 (B) shows the higher magnification micrograph of marked area 1 in Figure 4.7 (A). Figure 4.18: Scanning electron micrographs of an alloy with 12 wt. % Mn at: (A) 500X, (B) 1000X, and (C) 500X magnification. B C 1 A 56 The scanning electron micrographs of an alloy with 18 wt. % Mn in Figure 4.19 shows the carbide network along the grain boundaries. The marked area 2 in Figure 4.19 (A) and area 3 in Figure 4.19 (C) have been magnified in Figure 4.19 (B) and 4.19 (D), respectively. At few sites in Figure 4.19 (A), (B), (C), and (D), the grain boundaries of this alloy were decorated by the cluster of carbides. Figure 4.19: Scanning electron micrographs of an alloy with 18 wt. % Mn at: (A) 500X, (B) 1000X, (C) 500X, and (D) 500X magnification. A B C D 3 2 57 The carbides formed in an alloy with 23 wt. % Mn are shown in Figures 4.20 (A) and (C). The carbides appeared not only at the grain boundaries, but also inside the austenite grains. The marked area 4 in Figure 4.20 (A) can be seen clearly in Figure 4.20 (B). The carbide agglomeration networks on the grain boundaries were discontinuous and carbides were present either inside the grains or on the grain boundaries. Figure 4.20: Scanning electron micrographs of an alloy with 23 wt. % Mn at: (A) 500X, (B) 2000X, and (C) 1000X magnification. A C 4 B 58 The SEM analysis showed that the grain boundary precipitates were present in all the alloys. Thick continuous grain boundary delineations, which are typically regarded to be detrimental to the performance of the Hadfield steel, were not present in any of these steels. All the castings were heat-treated at 1065 ℃, which for most high Mn steels is considered adequate for dissolving carbides in austenite Fe matrix. For Hadfield steels containing Cr content, a higher solution heat-treatment is required due to the high thermal stability of the Cr carbides [37]. The formation of the grain boundary carbides could be attributed to the incomplete dissolution of the chromium carbides. X-ray Energy Dispersive Spectroscopy 4.2.3.1 EDS Map Scan of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn The EDS map of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn are shown in Figures 4.21, 4.22, and 4.23, respectively. The scans revealed that the distribution of C and P were uniform in all alloys whereas the distribution of Fe, Mn, and Cr locally varied. The distribution of Fe, Mn, and Cr along the sites of carbides suggests the possible formation of M3C and M23C6 –type carbides (M: metallic atom). The elemental distribution of C and Mn in the alloys suggested that the solution heat-treatment temperature was sufficient to homogenize the microstructure. Since a small volume fraction of carbides remained, the annealing temperature was not high enough for the complete dissolution of carbides such as M3C and M23C6, which often form in these steels [18]. The map scans for alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn reveal the dominant distribution of Cr and Mn (carbide forming elements) at the carbide sites, as shown in Figures 4.21-23, respectively. Gurol et al. [30] also observed the segregation of carbide forming elements such as Mn and Mo at the carbide sites in an alloy containing 17 wt. % Mn. 59 The Si distribution appeared in the map scan of an alloy with 12 wt. % Mn at the carbide’s site as shown in Figure 4.10. Additional EDS map scans of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn can be found in Appendix B. Mn Fe Cr Si C Figure 4.21: EDS map scan of an alloy with 12 wt. % Mn. 60 Mn Fe Cr Si C Figure 4.22: EDS map scan of an alloy with 18 wt. % Mn. 61 Mn Fe Cr Si C Figure 4.23: EDS map scan of an alloy with 23 wt. % Mn. 62 4.2.3.2 EDS Line Scan of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn The EDS line scans of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn are shown in Figures 4.24, 4.25, and 4.26 respectively. In all the alloys, the concentration of Fe, Mn, and Cr varied along a line passing the carbides, as seen in Figures 4.24, 4.25, and 4.26. The counts of Mn and Cr increased on the carbide-rich areas marked in Figure 4.24-4.26, while decreasing for Fe. The concentrations for other elements remained constant along the scan line in all alloys. Figure 4.24: EDS line scan of an alloy with 12 wt. % Mn. Mn Carbide rich area Fe Cr Mn 63 Mn Mn Figure 4.25: EDS line scan of an alloy with 18 wt. % Mn. Carbide rich area Fe Cr Mn 64 4.2.3.3 EDS Point/Area Scan of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn The composition of different areas marked in Figures 4.27, 4.28, and 4.29 is included in Tables 4.6, 4.7, and 4.8 respectively. In Figure 4.27, the areas A1, A2, and A3 were rich in Cr as compared to areas A4 and A5, suggesting that these islands were representative of blocky carbides rich in Cr. Similar clusters of carbides were seen in alloys with 18 wt. % Mn and 23 wt. % Mn. The EDS point/area scans validate the distribution of Cr and Mn on the carbide Carbide rich area Figure 4.26: EDS line scan of an alloy with 23 wt. % Mn. Carbide rich areas Fe Mn Cr 65 sites in all the alloys, as seen in the EDS map scan in Section 4.2.3.2. Thus, in all the alloys, the wt. % of Cr and Mn increased gradually at the selected area on the carbides, whereas the concentration of Fe decreased. Area Fe C Mn Cr Si A1 54.58 7.07 22.72 15.64 - A2 54.83 7.29 22.21 15.66 - A3 54.06 7.09 23.85 15.01 - A4 78.36 6.36 12.21 2.36 0.71 A5 74.67 5.69 15.72 3.13 0.80 Area Fe C Mn Cr Si A1 73.61 8.02 19.23 2.3 0.56 A2 71.18 8.0 18.05 2.23 0.54 A3 47.29 8.08 30.63 12.44 A4 51.52 8.12 29.27 10.77 0.11 A5 52.59 7.73 28.25 10.23 0.12 A6 71.66 8.69 21.34 2.53 0.64 A7 48.86 8.49 30.62 11.89 - A8 44.48 8.72 29.84 11.65 - Table 4.6: Chemical composition (wt. %) in an alloy with 12 wt. % Mn. Table 4.7: Chemical composition (wt. % ) in an alloy with 18 wt. % Mn. Figure 4.27: EDS point/area scan of an alloy with 12 wt. % Mn. A1 A2 A3 A5 A4 Figure 4.28: EDS point/area scan of an alloy with 18 wt. % Mn. A1 A2 A3 A4 A5 A7 A8 A6 66 4.3 Phase Analysis The phases of the alloys were studied using the X-ray diffraction technique. Figures 4.30, 4.31, and 4.32 show the XRD spectra of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn, respectively. Austenitic Fe and SiC were identified in all the alloys. Mn is a weak carbide forming element and doesn’t generate separate carbides in steels, but replaces the Fe in cementite to form (Fe,Mn)3C [51]. Fe1.8Mn1.2C carbides were detected in alloys with 12 wt. % Mn and 18 wt. % Mn. On the other hand, Fe1.1Mn3.9C2 was detected in an alloy with 23 wt. % Mn. This could be due to high Mn and C wt. % in alloy with 23 wt. % Mn. Cr is a strong carbide forming element and forms (Fe,Cr)3C, (Fe,Cr)7C6 and (Fe,Cr)23C6 type carbides in Fe-based alloys [51]. With Cr content up to 1.8% in Hadfield steels, (Fe,Cr)3C type of carbides were seen [4]. For Cr greater than 1.8 wt.%, the (Fe,Cr)23C6 type carbides were expected to Area Fe C Mn Cr Si A1 66.92 5.86 24.23 2.21 0.78 A2 67.99 5.29 23.76 2.18 0.79 A3 69.95 5.21 22.08 2.04 0.72 A4 45.10 7.81 34.15 12.81 0.13 A5 49.91 7.91 32.47 9.41 0.29 A6 57.14 8.23 28.6 5.53 0.5 A7 45.8 7.14 34.36 12.54 0.14 A8 43.04 7.42 34.66 13.22 0.12 Table 4.8: Chemical composition (wt. % ) in an alloy with 23 wt. % Mn. A2 A5 A6 A4 A7 A8 A3 A1 Figure 4.29: EDS point/area scan of an alloy with 23 wt. % Mn. A2 67 form [18]. Therefore, (Fe,Cr)23C6 was observed in all the alloys with slight variation in the wt.% of Fe and Cr. Figure 4.31: XRD pattern of an alloy with 18 wt. % Mn. Alloy with 18 wt. % Mn γ: Austenitic Fe (Score 65) 𝛥: SiC (Score 41) ♦: C6Cr18.93Fe4.07 (Score 27) ♥:Fe1.8Mn1.2C(Score 12) Figure 4.30: XRD pattern of an alloy with 12 wt. % Mn. Alloy with 12 wt. % Mn γ: Austenitic Fe (Score 62) 𝛥 : SiC (Score 47) ♦: C6Cr18.93Fe4.07 (Score 31) ♥ :Fe1.8Mn1.2C(Score 12) 68 4.4 Mechanical Testing Charpy-impact Testing Figure 4.33 shows the results of Charpy-Impact testing of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn. The alloy with 23 wt. % Mn absorbed the highest amount of energy (119.27±15.6 J) before failure, followed by alloys with 12 wt. % Mn and 18 wt. % Mn, respectively. The embrittlement of Hadfield steels is related to the percentage of grain boundaries coverage with carbides [26], [52]. The full austenite structure with fine grain size and low area fraction of carbides (as shown in Table 4.9) as compared to others may have contributed to a higher toughness (>100 J) in an alloy with 12 wt. % Mn. For an alloy with 18 wt. % Mn, the area fraction of carbides after the heat-treatment was highest and the carbide Alloy with 23 wt. % Mn γ: Austenitic Fe (Score 46) 𝛥: SiC (Score 32) ♦: C6Cr15.58Fe7.42 (Score 25) ♥: Fe1.1Mn3.9C2 (Score 15) Figure 4.32: XRD pattern of an alloy with 23 wt. % Mn. 69 network had covered more than 50% of the grain boundaries at several sites (as seen in Figure 4.2). Therefore, this could be one of the reasons for the lowest Charpy-impact value of an alloy with 18 wt. % Mn (resulting in a brittle material fracture due to the presence of carbides [4], [7]). Although an alloy with 23 wt. % Mn had a high average grain size and high area fraction of carbides in comparison with an alloy with 12 wt. % Mn, it resulted in a higher Charpy-impact value. This could be due to solid solution strengthening due to higher C and Mn content retained in the austenitic Fe matrix after the heat treatment [4]. Table 4.9: Summary of microstructure analysis and wt.% of Mn and C in alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn addition. Alloy Wt.% of Mn Wt.% of C Average Grain Size (µm) Area Fraction of Carbides (%) 12 wt. % Mn 12.14 1.05 116.6 2.2 ± 0.8 18 wt. % Mn 17.70 1.13 175.6 5.0 ± 0.7 23 wt. % Mn 22.78 1.15 445.6 4.2 ± 0.9 Figure 4.33: Charpy-impact values of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn. 119.27 ± 15.665.03 ± 1.5 124.70 ± 5.4 02040608010012014012 wt. % Mn 18 wt. % Mn 23 wt. % MnCharpy Impact (Joules)Alloy70 Tensile Testing Figure 4.34 shows the results of tensile testing. With an increase in the manganese content, the ductility and ultimate tensile strength decreased, whereas the yield strength increased. The percentage of carbon in an alloy with 12 wt. % Mn was 1.05%, whereas, for alloys with 18 wt. % Mn and 23 wt. % Mn it was around 1.13 and 1.15 wt. % C, respectively. The increase in the percentage of C and Mn can be one of the reasons for the increase in the yield strength due to solid solution strengthening [3]. However, with an increase in wt. % of C and Mn, the ultimate tensile strength, and elongation decreased due to undissolved carbides (as measured by phase area fraction) in the microstructure after heat treatment [4], [18]. The austenite microstructure of an alloy with 12 wt. % Mn with fine grain size and low area fraction of carbides (as shown in Table 4.9) resulted in the highest ultimate tensile strength and high ductility. Figure 4.34: Tensile testing results of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn. a 71 Microhardness Figure 4.35 shows the average initial microhardness of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn. With an increase in Mn and C wt. % in Hadfield steel, the increase in hardness was observed due to retained C and Mn in a solid solution in austenitic Fe matrix [21], [30]. It was also observed that the hardness of Hadfield steels increases due to carbides precipitation in the microstructure [4], [6], [7]. The area fraction of carbides in an alloy with 12 wt. Mn % was lowest as compared to alloys with 18 wt. % Mn, and 23 wt. % Mn, which may result in the lowest average hardness for an alloy with 12 wt. % Mn. Although the area fraction of carbides, as previously shown in Table 4.9, was highest in an alloy with 18 wt.% Mn, the average hardness value of an alloy with 23 wt. % Mn was higher. This can be due to the uniform distribution of carbides on the grain boundaries and within the grains, as seen in Figure 4.14, which improved the overall hardness of an alloy with 23 wt. % Mn. Figure 4.35: Average hardness of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn. 268±3.1278±2.8281±3.326526726927127327527727928128312 wt. % Mn 18 wt. % Mn 23 wt. % MnVicker Hardness (HV)Alloy72 Figure 4.36 shows the measured microhardness value of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn during the wear test at different pseudo-wear testing periods (time = 0, 12, 72, 132, and 192 minutes). The SiC paper was changed after every 1 minute for an initial 12 min then changed after every 5 min for the rest of the testing periods. The microhardness of all the alloys slightly increased in the initial 12 minutes of testing. Post 12 min, it remained almost constant in all the alloys for the next three 60 minutes testing periods. It was observed that the increase in hardness of the Hadfield steels was not significant during the abrasion, as there was no plastic deformation [18]. The rate of work hardening of this steel is often directly related to the magnitude of high impact loading [53]. The hardness value above 500 HV has been reported on the work-hardened surface of Hadfield steels [52]. The hardness value increases during the impact loading, which improves the wear resistance of Hadfield steels [54]. Therefore, severe plastic deformation is the main key for the work hardenability of Hadfield steel which improved the wear resistance. Figure 4.36: Average hardness value of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn during the pseudo-wear test. 73 Pseudo-wear Testing Figure 4.37 shows the weight percentage loss in alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn after the 192 minutes of wear testing. The alloys with 18 wt. % Mn and 23 wt. % Mn showed improved resistance against wear, with approximately 8.1% weight loss. The weight loss % was maximum around 8.6 % in an alloy with 12 wt. % Mn. The hardness plays a major role in the wear resistance of Hadfield steels. With an increase in wt.% of Mn in Hadfield steels, the hardness increased due to solid solution strengthening after the heat-treatment in austenitic Fe [7]. The formation of undissolved carbides after the heat-treatment increased with an increase in C wt.% [3,10]. The phase area fraction of the carbides in the microstructure of alloys with 18 wt. % Mn and 23 wt. % Mn was higher as compared to an alloy with 12 wt. % Mn as shown in Table 4.9. The high fraction of carbides may be the reason for the good wear resistance of alloys with 18 wt. % Mn and 23 wt. % Mn compared to 12 wt. % Mn alloy [2]. 01234567891 3 5 7 9 11 17 27 37 47 57 67 77 87 97 107117127137147157167177187Weight Loss (%)Testing Time (Minutes)12 wt. % Mn 18 wt. % Mn 23 wt. % MnFigure 4.37: Weight loss percentage of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn Vs time. 74 Chapter 5: Conclusions and Recommended Future Work This chapter presents a summary of this thesis work. In this work, three different amounts of Mn wt. % (approximately 12, 18, and 23) were added to the Hadfield steels through the casting process. Further, the specimens were characterized based on their microstructure, phases, and mechanical properties. The conclusion is divided into two sections: (1) Fe-Mn-C System, Microstructure, and Phase Analysis (2) Mechanical Properties. 5.1 Fe-Mn-C System, Microstructure, and Phase Analysis The liquidus projection of the Fe-Mn-C system revealed the decrease in liquidus temperature with an increase in wt. % of Mn from 12.14 to 22.78 wt. %. The results of optical and scanning electron microscopy revealed the presence of intermittent carbides on the grain boundaries in all the alloys. The intergranular carbides were present in alloys with 18 wt. % Mn and 23 wt. % Mn. The formation of the envelope of carbides covering more than 50% grain boundaries was observed in an alloy with 18 wt. % Mn. The area fraction of carbides was around 5 % for an alloy with 18 wt. % Mn, followed by 4.2% and 2.2% for alloys with 23 wt. % Mn and 12 wt. % Mn, respectively. The EDS map scan revealed the homogenous distribution of elements in the SEM micrographs after the heat-treatment. The accumulation of Cr and Mn was observed in all the alloys at the carbide’s sites during the EDS line and point/area scan. XRD analysis found out that phases such as austenite Fe, SiC, and (Fe,Cr)23C6 were present in all the alloys with a slight variation of Cr and Fe in (Fe,Cr)23C6. The Fe1.8Mn1.2C was present in alloys with 12 wt. % Mn and 18 wt. % Mn whereas Fe1.1Mn3.9C2 was present in an alloy with 23 wt. % Mn. 75 5.2 Mechanical Properties An alloy with 23 wt. % Mn showed the best resistance against the high strain impact in the Charpy-impact testing and absorbed 124.7 J of energy. The alloy with 12 wt. % Mn and 18 wt. % Mn absorbed 119.27 Joules and 65.03 J, respectively. The yield strength was highest in an alloy with 18 wt. % Mn at 450 MPa followed by alloys with 23 wt. % Mn and 12 wt. % Mn at 442 and 399 MPa, respectively. The ultimate tensile strength and elongation decreased with an increase in the wt.% of Mn in the alloys. The wt. loss % in alloys with 18 wt. % Mn, and 23 wt. % Mn after 192 minutes of abrasion in pseudo-wear testing was around 8.1%. The weight loss % was slightly higher in an alloy with 12 wt. % Mn. The average hardness value of alloys with 12 wt. % Mn, 18 wt. % Mn, and 23 wt. % Mn was 268, 278, and 282 HV, respectively. This implied that the average hardness value increased with an increase in wt.% of Mn in the alloys. The average hardness value of the alloys was also measured at the pseudo testing time: 12, 72, 132, and 192 minutes. After the initial 12 minutes of pseudo-wear testing, a slight increase in the average hardness value of all the alloys was observed. 5.3 Recommendations for Future Work This section provides a few recommendations for future work that can be done on the Hadfield steels. Low-Temperature Charpy-impact and Tensile Testing The temperature at the mining sites in the province of Saskatchewan and Ontario remains below 0 ℃ for more than 30 % of the year. Conducting the low temperature (-50,-30, and 0℃) Charpy-impact and tensile testing will help in evaluating the behavior of Hadfield steels with varying Mn wt. % in low-temperature conditions. 76 Transmission Electron Microscope analysis Different phases with different compositions were revealed in the XRD analysis. Conducting the transmission electron microscope (TEM) analysis would help to evaluate the lattice mismatch between the carbides and austenitic Fe matrix in Hadfield steels with varying Mn wt. %. Impact Wear Testing The weight loss % of the alloys was evaluated by abrading the alloys on the SiC papers. The impact wear due to rocks on the crusher’s surface also happens in the actual crushing process. Conducing the Impact Wear Testing will help to explore the effect of varying Mn wt. % on the impact wear resistance of Hadfield Steels. 77 Bibliography [1] S. K. Haldar, Mineral Processing. 2018. [2] J. O. Agunsoye, S. I. Talabi, and O. Bello, “Wear characteristics of heat-treated hadfield austenitic manganese steel for engineering application,” Adv. Prod. Eng. Manag., vol. 10, no. 2, pp. 97–107, 2015, doi: 10.14743/apem2015.2.195. [3] J. (Jimmy) Jiang, Y. Xie, W. Qian, and R. Hall, “Important factors affecting the gouging abrasion resistance of materials,” Miner. Eng., vol. 128, no. May, pp. 238–246, 2018, doi: 10.1016/j.mineng.2018.09.001. [4] D. K. Subramanyam, A. E. Swansiger, and H. S. Avery, “Austenitic manganese steels,” Properties and Selection: Irons, Steels, and High-Performance Alloys, vol. 1. ASM International, p. 0, Jan. 01, 1990, doi: 10.31399/asm.hb.v01.a0001045. [5] A. Studnicki, J. Kilarski, M. Przybył, J. Suchoń, and D. Bartocha, “Wear resistance of chromium cast iron – research and application,” Manuf. Eng., vol. 16, no. 1, pp. 63–73, 2006. [6] R. Lencina, C. Caletti, K. Brunelli, and R. Micone, “Assessing wear performance of two high-carbon Hadfield steels through field tests in the mining industry,” Procedia Mater. Sci., vol. 9, pp. 358–366, 2015, doi: 10.1016/j.mspro.2015.05.005. [7] S. A. Torabi, K. Amini, and M. Naseri, “Investigating the effect of manganese content on the properties of high manganese austenitic steels,” Int. J. Adv. Des. Manuf. Technol., vol. 10, no. 1, pp. 75–83, 2017. [8] A. K. Srivastava and K. Das, “Microstructural characterization of Hadfield austenitic manganese steel,” J. Mater. Sci., vol. 43, no. 16, pp. 5654–5658, 2008, doi: 78 10.1007/s10853-008-2759-y. [9] M. Ericsson and O. Löf, “Mining’s contribution to national economies between 1996 and 2016,” Miner. Econ., pp. 223–250, 2019, doi: 10.1007/s13563-019-00191-6. [10] B. Marshall, “The state of Canada’s mining industry,” 2019, [Online]. Available: https://mining.ca/documents/facts-and-figures-2019/. [11] I. The, J. C. Jaw, T. Blake, and B. Crusher, “Jaw crusher,” Pit Quarr., vol. 95, no. 1, p. 24, 2002, doi: 10.1016/b978-044451636-7/50005-x. [12] D. Li, “Abrasive wear,” Friction, Lubrication, and Wear Technology, vol. 18. ASM International, p. 0, Dec. 31, 2017, doi: 10.31399/asm.hb.v18.a0006382. [13] J. Tasker and M. O. H. Amuda, “Austenitic steels: non-stainless,” Ref. Modul. Mater. Sci. Mater. Eng., vol. 93, no. November 2015, pp. 1–7, 2017, doi: 10.1016/b978-0-12-803581-8.09206-7. [14] C. S. Mahlami and X. Pan, “Challenges and developments of hadfield manganese steel castings based on service life,” in Proceedings of the 71th World Foundry Congress, Bilbao, Spain, 2014, pp. 19–21. [15] Matecsolutions.com. 2022. hi manganese chrome-Matec Solutions. [online] Available at: https://www.matecsolutions.com/js. [16] P. Chowdhury, D. Canadinc, and H. Sehitoglu, “On deformation behavior of Fe-Mn based structural alloys,” Mater. Sci. Eng. R Reports, vol. 122, pp. 1–28, 2017, doi: 10.1016/j.mser.2017.09.002. [17] D. Havel, “Austenitic manganese steel a complete overview,” SFSA Tech. Oper., pp. 1–19, 2017. [18] M. Sabzi and M. Farzam, “Hadfield manganese austenitic steel: A review of 79 manufacturing processes and properties,” Mater. Res. Express, vol. 6, no. 10, 2019, doi: 10.1088/2053-1591/ab3ee3. [19] D. V. Lychagin, A. V. Filippov, O. S. Novitskaia, Y. I. Chumlyakov, E. A. Kolubaev, and O. V. Sizova, “Friction-induced slip band relief of -Hadfield steel single crystal oriented for multiple slip deformation,” Wear, vol. 374–375, pp. 5–14, 2017, doi: 10.1016/j.wear.2016.12.028. [20] Ma, L., Huang, C., Dolman, K., Tang, X., Yang, J., Shi, Z. and Liu, Z.S., “A method to calculate the bulk hardness of metal matrix composite using Hadfield steel reinforced with niobium carbide particles as an example,” Mech. Mater., vol. 112, pp. 154–162, 2017, doi: 10.1016/j.mechmat.2017.06.005. [21] D. K. Subramanyam and H. S. Avery, “Properties and selection: irons, steels, and high performance alloys,” ASM handbook, ASM Int., 1999. [22] ASTM A128/A128M, “Standard specification for steel castings, austentic manganese,” Quality, vol. 93, no. Reapproved 2012, pp. 1–3, 2003, doi: 10.1520/A0128. [23] H. Okamoto, “Fe (iron) binary alloy phase diagrams,” Alloy Phase Diagrams, vol. 25, no. 1, pp. 340–361, 2018, doi: 10.31399/asm.hb.v03.a0006162. [24] X. Liang, F. Haakonsen, J. I. Erauskin, A. Sargyan, and J. L. Arana, Optimizing of Strømhard austenitic manganese steel, vol. 49, no. 4. 2009. [25] S. A. Balogun, D. E. Esezobor, and J. O. Agunsoye, “Effect of melting temperature on the wear characteristics of austenitic manganese steel,” J. Miner. Mater. Charact. Eng., vol. 07, no. 03, pp. 277–289, 2008, doi: 10.4236/jmmce.2008.73021. [26] R. Zavadil and S. Kuyucak, “Microstructure vs. impact toughness relationship in Hadfield’s austenitic manganese steel,” Microsc. Microanal., vol. 8, no. SUPPL. 2, pp. 80 1290–1291, 2002, doi: 10.1017/s1431927602104880. [27] M. Beheshti, M. Zabihiazadboni, M. C. Ismail, S. Kakooei, and S. Shahrestani, “Investigation on simultaneous effects of shot peen and austenitizing time and temperature on grain size and microstructure of austenitic manganese steel (Hadfield),” IOP Conf. Ser. Mater. Sci. Eng., vol. 328, no. 1, 2018, doi: 10.1088/1757-899X/328/1/012006. [28] C. Chen, F. C. Zhang, F. Wang, H. Liu, and B. D. Yu, “Effect of N+Cr alloying on the microstructures and tensile properties of Hadfield steel,” Mater. Sci. Eng. A, vol. 679, no. June 2016, pp. 95–103, 2017, doi: 10.1016/j.msea.2016.09.106. [29] C. L. Projection, “C (carbon) ternary alloy phase diagrams,” Alloy Phase Diagrams, vol. 26, pp. 683–702, 2018, doi: 10.31399/asm.hb.v03.a0006236. [30] U. Gürol and S. C. Kurnaz, “Effect of carbon and manganese content on the microstructure and mechanical properties of high manganese austenitic steel,” J. Min. Metall. Sect. B Metall., vol. 56, no. 2, pp. 171–182, 2020, doi: 10.2298/JMMB191111009G. [31] V. Lazić, D. Milosavljević, S. Aleksandrović, P. Marinković, G. Bogdanović, and B. Nedeljković, “Carbide type influence on tribological properties of hard faced steel layer - part I - theoretical considerations,” Tribol. Ind., vol. 32, no. 2, pp. 11–20, 2010. [32] M. Sabzi, S. M. Far, and S. M. Dezfuli, “Effect of melting temperature on microstructural evolutions, behavior and corrosion morphology of Hadfield austenitic manganese steel in the casting process,” Int. J. Miner. Metall. Mater., vol. 25, no. 12, pp. 1431–1438, 2018, doi: 10.1007/s12613-018-1697-1. [33] H. R. Jafarian, M. Sabzi, S. H. Mousavi Anijdan, A. R. Eivani, and N. Park, “The 81 influence of austenitization temperature on microstructural developments, mechanical properties, fracture mode and wear mechanism of Hadfield high manganese steel,” Journal of Materials Research and Technology, vol. 10. pp. 819–831, 2021, doi: 10.1016/j.jmrt.2020.12.003. [34] Michalon, D., G. Mazet, and Ch Burgio. "Manganese steel for abrasive environments: A conditioning process for Hadfield's manganese steel and a novel method of producing FAM bearings from the same material." TRIBOLOGY international 9.4 (1976): 171-178. [35] M. Azadi, A. M. Pazuki, and M. J. Olya, “The effect of new double solution heat treatment on the high manganese Hadfield steel properties,” Metallogr. Microstruct. Anal., vol. 7, no. 5, pp. 618–626, 2018, doi: 10.1007/s13632-018-0471-0. [36] G. Tęcza and S. Sobula, “Effect of heat treatment on change microstructure of cast high-manganese Hadfield steel with elevated chromium content,” Arch. Foundry Eng., vol. 14, pp. 67–70, 2014. [37] S. Kuyucak, “Austenitic manganese steel castings,” Mater. Park. OH ASM Int. 2004., pp. 701–708, 2004. [38] Y. N. Dastur and W. C. Leslie, “Mechanism of work hardening in Hadfield manganese steel,” Metall. Trans. A, vol. 12, no. 5, pp. 749–759, 1981. [39] Raghavan KS, Sastri AS, and Marcinkowski MJ, “Nature of work-hardening behavior in Hadfield’s manganese steel,” Met Soc AIME-Trans, vol. 245, no. 7, pp. 1569–1575, 1969. [40] W. S. Owen and M. Grujicic, “Strain aging of austenitic Hadfield manganese steel,” Acta Mater., vol. 47, no. 1, pp. 111–126, 1998, doi: 10.1016/S1359-6454(98)00347-4. 82 [41] I. Karaman, H. Sehitoglu, A. J. Beaudoev, Y. I. Chumlyakov, H. J. Maier, and C. N. Tomé, “Modeling the deformation behavior of Hadfield steel single and polycrystals due to twinning and slip,” Acta Mater., vol. 48, no. 9, pp. 2031–2047, 2000, doi: 10.1016/S1359-6454(00)00051-3. [42] I. Karaman, H. Sehitoglu, K. Gall, Y. I. Chumlyakov, and H. J. Maier, “Deformation of single crystal Hadfield steel by twinning and slip,” Acta Mater., vol. 48, no. 6, pp. 1345–1359, 2000, doi: 10.1016/S1359-6454(99)00383-3. [43] O. O. Lambakakhar and Y. I. Paskal, “Dislocation structure and work-hardening of manganese steel,” Sov. Phys. J., vol. 16, no. 7, pp. 910–914, 1973. [44] W. Yan, L. Fang, K. Sun, and Y. Xu, “Effect of surface work hardening on wear behavior of Hadfield steel,” Mater. Sci. Eng. A, vol. 460–461, pp. 542–549, 2007, doi: 10.1016/j.msea.2007.02.094. [45] ASTM E112-10, “Standard test methods for determining average grain size,” Standards, pp. 1–27, 2010, doi: 10.1520/E0112-13.1.4. [46] ASTM E 23-12c, “Standard test methods for notched bar impact testing of metallic materials,” Standards, pp. 1–25, 2012, doi: 10.1520/E0023-18. [47] ASTM E8, “ASTM E8/E8M standard test methods for tension testing of metallic materials 1,” Annu. B. ASTM Stand. 4, no. C, pp. 1–27, 2010, doi: 10.1520/E0008. [48] K. Hardness, “Standard test methods for vickers hardness and knoop hardness of metallic materials bt - standard test methods for vickers hardness and knoop hardness of metallic materials,” vol. i, pp. 1–27, 17AD, doi: 10.1520/E0092-17.2. [49] V. Průcha, Z. Jansa, J. Šimeček, O. Žďánský, and A. Kříž, “Characterization of microstructure of Hadfield steel,” Solid State Phenom., vol. 270 SSP, pp. 265–270, 83 2017, doi: 10.4028/www.scientific.net/SSP.270.265. [50] J. O. Olawale, S. A. Ibitoye, and M. D. Shittu, “Workhardening behaviour and microstructural analysis of failed austenitic manganese steel crusher jaws,” Mater. Res., vol. 16, no. 6, pp. 1274–1281, 2013, doi: 10.1590/S1516-14392013005000144. [51] V. Lazić, D. Milosavljević, S. Aleksandrović, P. Marinković, G. Bogdanović, and B. Nedeljković, “Carbide type influence on tribological properties of hard faced steel layer–part I–theoretical considerations,” Tribol. Ind., vol. 32, no. 2, pp. 11–20, 2010. [52] C. S. Mahlami and X. Pan, “An overview on high manganese steel casting,” World Foundry Congrees, pp. 1–10, 2014. [53] Y. H. Wen, H. B. Peng, H. T. Si, R. L. Xiong, and D. Raabe, “A novel high manganese austenitic steel with higher work hardening capacity and much lower impact deformation than Hadfield manganese steel,” Mater. Des., vol. 55, pp. 798–804, 2014, doi: 10.1016/j.matdes.2013.09.057. [54] Y. H. Xu, L. Fang, Q. Cen, and J. H. Zhu, “Subsurface microstructure evolution of hadfield steel under high impact energy,” Mater. Sci. Forum, vol. 475–479, pp. 117–120, 2005, doi: 10.4028/www.scientific.net/msf.475-479.117. 84 Appendices Appendix A A B C D Figure A.1: Optical micrographs of an alloy with 12 wt. % Mn at: (A) 50X, (B)100X , (C)500X, and (D) 500X magnification. Figure A.2: Optical micrographs of an alloy with 18 wt. % Mn at: (A) 50X, (B)100X , (C)500X, and (D) 500X magnification. A B D C 85 Figure A.3: Optical micrographs of an alloy with 23 wt. % Mn at: (A) 50X, (B)100X , (C)500X, and (D) 500X magnification. A B C D 86 Appendix B Figure B.1: EDS map scan of an alloy with 12 wt. % Mn. Mn Cr Si C Fe 87 Fe Mn Si Cr C Figure B.2: EDS map scan of an alloy with 18 wt. % Mn. 88 Fe Mn Cr Si C Figure B.3: EDS map scan of an alloy with 23 wt. % Mn.
Thesis/Dissertation
eng
Mechanical Engineering
Vancouver : University of British Columbia Library
Attribution-NonCommercial-NoDerivatives 4.0 International
Graduate
Bichler, Lukas
Investigating the effect of Mn content on the microstructure and mechanical properties of high manganese steel
Text
Still Image