Development of Mg-Al-Sn and Mg-Al-Sn-Si Alloys and Optimization of Super Vacuum

Die Casting Process for Lightweight Applications

DISSERTATION

Presented in Partial Fulfillment of the Requirements for the Degree Doctor of Philosophy in the Graduate School of The Ohio State University

Andrew Daniel Klarner

Graduate Program in Materials Science and Engineering

The Ohio State University

2018

Dissertation Committee:

Dr. Alan A. Luo, Advisor

Dr. Glenn S. Daehn

Dr. Michael J. Mills

Dr. Gary Kennedy

Copyrighted by

Andrew Daniel Klarner

2018

Abstract

Light weighting of structural components is crucial to lessen environmental impact, lower costs, and to reduce energy consumption. The most popular method to accomplish this is to replace materials, such as steel, with less dense materials such as aluminum or magnesium. However these lighter materials cannot always meet the strength requirements of structural applications, therefore further material development is needed to improve their mechanical properties. This dissertation will summarize the design and development of two new magnesium alloys for use in the high pressure die casting process (HPDC); Mg-Al-Sn (AT) and Mg-Al-Sn-Si (ATS). These alloys were designed to improve upon current HPDC magnesium alloys AM50/60 (Mg-5/6wt.% Al-

0.2wt.% Mn) and AZ91 (Mg-9wt.% Al-1wt.% Zn), which possess limited mechanical properties. The CALPHAD (CALculation of PHAse Diagrams) method was used in the development of AT and ATS alloys and to aid in the design of heat treatment schedules.

With this method thermodynamic models and limited thermodynamic data are combined to produce predictions of phase equilibria, solidification sequences, and other useful information about multicomponent alloy systems. This allows for a more streamlined and cost effective alloy development process that can be used to design tailored alloy systems. The AT and ATS alloys were evaluated with scanning electron microscopy

(SEM), energy dispersive x-ray spectroscopy (EDS), high-angle annular dark field scanning transmission electron microscopy (HAADF-STEM), and transmission electron microscopy (TEM) to characterize the microstructure of the alloys in the as-cast condition as well as in multiple heat treated conditions. Mechanical testing was performed on specimens that were produced by the HPDC process to compare the strength and ductility of these alloys to currently used magnesium alloys.

The 2nd part of this dissertation explores the development and optimization of super vacuum die casting process for aluminum and magnesium thin-wall castings, using process simulation and experimental validation. High pressure die casting is a processing technique that produces metal components by forcing molten metal into steel molds, called dies, at very high speeds and pressures. There are many advantages to using this process including the ability to produce complex near net shape parts at very high production rates with high repeatability. However because of the need to produce very large castings which have thin wall and also to limit the porosity in the casting, further process development is needed. Super Vacuum Die Casting (SVDC) has been developed recently and is a technique that introduces a vacuum in the die cavity before filling. With the absence of air in the die cavity, the porosity found in castings can be limited and the filling of castings improved. Two experimental dies, i.e., test specimen die and fluidity die, were designed to evaluate the castability of several new Al and Mg alloys and optimize process parameters for these alloys. The process conditions were successfully validated in industrial castings such as a door inner and a side impact beam.

iii

Dedicated to my Parents

Acknowledgments

I would first like to acknowledge my advisor, Dr. Alan Luo, for his continued support and guidance throughout my time in graduate school. He provided me with many research opportunities to expand my knowledge in alloy development and materials manufacturing. I would like to thank all my colleagues in the Light Metals and

Manufacturing Research Laboratory, especially Dr. Weihua Sun and Janet Meier, for their advice and assistance with my research. Bill Tullos and Dr. Jerry Brevick, both of the Integrated Systems Engineering department, were essential in the manufacturing of experimental dies and in operation of the die cast machine, I sincerely thank them for their support and guidance. I would also like to acknowledge the Simulation Innovation and Modeling Center (SIMCenter) and the Center for Design and Manufacturing

Excellence (CDME) for providing me with the necessary tools to successfully conduct my research. Dr. Jon Carter of General Motors provided support in many aspects which improved the outcome of this research; I am appreciative of his assistance. I would like to thank the Melt5B team, especially the leadership of Mike Zolnowski, for providing great insight in all things related to high pressure die casting as well as Light Weight

Innovations for Tomorrow (LIFT) for providing financial support.

In addition, I would like to thank my dissertation committee members, Dr.

Michael Mills and Dr. Glenn Daehn for their constructive suggestions and support which made this work more complete. My parents and brothers who supported me throughout my studies and provided constant encouragement for me to achieve my goals, I am truly grateful to have them in my life. I would like to shout out Leadbelly the OSU Ultimate

Team, as well as the Columbus Ultimate Community, who kept me physically and mentally young during my time in graduate school. And lastly to all the friends I met during and before Graduate school, I am thankful for your support and to have all of you involved in my life.

Vita

2008...... Farragut High School, Knoxville TN

2012...... B.S. Department of Materials Science and

Engineering, The University of Tennessee

2015...... M.S. Department of Materials Science and

Engineering, The Ohio State University

2012 to present ...... Graduate Research Associate, Light Metals

and Manufacturing Research Laboratory,

Department of Materials Science and

Engineering, The Ohio State University

Publications

1. A.D. Klarner, W. Sun, J. Meier, A.A. Luo, “Development of Mg-Al-Sn-Si

Alloys Using a Calphad Approach”, in Magnesium Technology, The Minerals, Metals &

Materials Society, Nashville, TN, 2016, pp. 79-82.

vii

2. A.D. Klarner, W. Sun, J. Miao, A.A. Luo, “Microstructure and Mechanical

Properties of High Pressure Die Cast Mg-Al-Sn-Si Alloys”, in Magnesium Technology,

The Minerals, Metals & Materials Society, San Diego, CA, 2017, pp. 289-295.

3. A.D. Klarner, E. Cinkilic, Y. Lu, J. Brevick, A.A. Luo, J. Shah, M. Zolnowski,

X. Yan, “A New Fluidity Die for Castability Evaluation of High Pressure Die Cast

Alloys”, NADCA Die Casting Congress and Tabletop, NADCA, Atlanta, GA, 2017

4. H. Ibrahim, A.D. Klarner, B. Poorganji, D. Dean, A.A. Luo and M. Elahinia,

"Microstructural, Mechanical and Corrosion Characteristics of Heat-Treated Mg-1.2Zn-

0.5Ca (wt.%) Alloy for Use as Resorbable Bone Fixation Material", Journal of the

Mechanical Behavior of Biomedical Materials, Volume 69, 2017, pp. 203-212.

Fields of Study

Major Field: Materials Science and Engineering

viii

Table of Contents

Abstract ...... ii

Acknowledgments...... v

Vita ...... vii

Publications ...... vii

Fields of Study ...... viii

Table of Contents ...... ix

List of Tables ...... xii

List of Figures ...... xiv

Chapter 1: Introduction ...... 1

Chapter 2: Background ...... 4

2.1: Properties of Magnesium ...... 4

2.1.1: Strengthening Mechanisms ...... 7

2.1.2: Magnesium Alloys...... 12

2.2: CALPHAD ...... 17

2.3: High Pressure Die Casting ...... 20 ix

2.3.1: Cold Chamber Die Casting ...... 23

2.3.2: Hot Chamber Die Casting ...... 24

2.3.3: Molten Magnesium Processing ...... 26

2.3.4: Molten Aluminum Processing ...... 28

2.4: Super Vacuum Die Casting ...... 30

Chapter 3: Mg-Al-Sn Alloy Development ...... 36

3.1: Introduction ...... 36

3.2: Sample Preparation and Methods...... 39

3.3: HPDC AT72 As-Cast Microstructure ...... 42

3.4: Heat Treatment Design and Microstructure ...... 50

3.5: HPDC AT72 As-Cast, T5, and T6 Mechanical Properties ...... 59

3.6: Conclusions ...... 66

Chapter 4: Mg-Al-Sn-Si Alloy Development ...... 68

4.1: Introduction ...... 68

4.2: Gravity Cast Experimental Procedure ...... 69

4.3: Gravity Cast Microstructure and Hardness ...... 71

4.4: Gravity Cast T5 and T6 Microstructure and Hardness ...... 76

4.5: HPDC ATS Sample Preparation and Methods ...... 82

4.6: HPDC As-Cast Microstructure and Mechanical Properties ...... 85

4.7: HPDC Heat Treatment Design ...... 94

4.8: HPDC Heat Treated Microstructure and Mechanical Properties ...... 98

4.9: Conclusions ...... 113

Chapter 5: High Pressure Die Casting Process Development ...... 117

5.1: Fluidity Studies of Aluminum Alloys ...... 117

5.1.1: Introduction ...... 117

5.1.2: Die Design and Experimental Procedure ...... 119

5.1.3: Results and Discussion ...... 124

5.1.4: Conclusions ...... 139

5.2: Super Vacuum Die Casting Development ...... 141

5.2.1: Introduction ...... 141

5.2.2: Die Design and Experimental Procedure ...... 144

5.2.3: Results and Discussion ...... 149

5.2.4: CANMET industry application ...... 157

5.2.5: Conclusions ...... 163

Chapter 6: Summary and Future Work ...... 165

References ...... 170

List of Tables

Table 1: Solubility data (for binary Mg-x systems) for common magnesium alloying additions[19] ...... 13

Table 2: Compositions found with inductively couple plasma optical emission spectrometry (ICP-OES) of each magnesium alloy; AM60B, AZ91, AT72, and ATS.... 40

Table 3: Mechanical properties of AM60B, AZ91, and AT72 from tensile specimen cut from HPDC thin-wall castings ...... 61

Table 4: Mechanical properties of AT72 specimen cut from HPDC thin-wall castings in different heat treatment conditions ...... 63

Table 5 Nominal compositions (wt.%) of the gravity cast Mg-Al-Sn-Si alloys ...... 70

Table 6: As-cast hardness values (Rockwell 15T) for gravity cast ATS alloys and AM50

...... 73

Table 7: Hardness (Rockwell 15T) values of gravity cast ATS alloys in the as-cast, solution treated, and T6 peak aged condition ...... 80

Table 8: Hardness (Rockwell 15T) values of gravity cast ATS alloys for as-cast, peak T5, and max T5 hardness ...... 81

Table 9: Composition of ATS determined by Inductively Coupled Plasma Mass

Spectrometry (ICP–OES)...... 83

xii

Table 10: Rockwell 15T hardness for HPDC AT72 and ATS in the as-cast condition.... 90

Table 11: Mechanical properties of HPDC AT72 and ATS tensile bars ...... 91

Table 12: Rockwell 15T hardness for HPDC ATS in the as-cast and solution treated conditions ...... 99

Table 13: Mechanical properties of HPDC ATS tensile bars in multiple heat treated conditions ...... 111

Table 14: Compositions (wt.%) of the studied alloys obtained by optical emission spectrometry ...... 122

Table 15: The liquidus, solidus, and fraction of eutectic liquid for each experimental alloy found with Pandat and PanAl2017 thermodynamic database...... 125

Table 16: The average flow length for each wall-thickness for both the EZCast and

Lift380 alloy at similar superheats and die temperatures ...... 131

Table 17: HTC values used in the ProCAST simulation obtained from Guo et al. [99] 135

Table 18: X-Ray porosity results for tensile bars produced at different vacuum levels . 150

Table 19: X-Ray porosity results for 3mm plates produced at different vacuum levels 151

Table 20: Mechanical properties for Lift380 produced at different vacuum levels ...... 153

xiii

List of Figures

Figure 1: Schematic of the hcp crystal structure and lattice parameters of magnesium ..... 6

Figure 2: Schematic showing the (a) slip systems and (b) twin systems present in magnesium alloys [9] ...... 7

Figure 3: Mg-Al phase diagram calculated by Pandat ...... 15

Figure 4: Microstructure of HPDC AM50 [26] ...... 16

Figure 5: Schematic of the cold chamber die casting process [38]...... 24

Figure 6: Schematic of the hot chamber die casting process [38] ...... 25

Figure 7: Simple diagram showing the vacuum die casting process; (a) vacuum starts after piston has passed the pouring hole (b) and continues pulling vacuum until casting is filled [46] ...... 33

Figure 8: Examples of chill blocks used in vacuum die casting [47] ...... 34

Figure 9: Photograph showing the dimensions of the HPDC door panel castings and location of tensile bar extraction ...... 41

Figure 10: The predicted solidification path for Mg-7Al-2Sn using the Scheil-Gulliver model and PanMG17 thermodynamic database ...... 45

Figure 11: Optical micrograph showing the as-cast microstructure of HPDC AT72 ...... 46

Figure 12: Backscatter electron image of the as-cast microstructure of HPDC AT72 ..... 48

xiv

Figure 13: Predicted concentration profiles of both (a) Al and (b) Sn in primary (Mg) for the AT72 alloy based on the Scheil-Gulliver model ...... 49

Figure 14: The predicated equilibrium phase fraction vs temperature curve for Mg-7Al-

2Sn ...... 52

Figure 15: DICTRA simulation showing the homogenization of (a) Al and (b) Sn in primary (Mg) at 420°C held for different hold times ...... 53

Figure 16: Optical micrograph showing the solution treated microstructure of HPDC

AT72 ...... 54

Figure 17: Backscatter electron image of the microstructure of solution treated (420°C for

8 hours) HPDC AT72 ...... 55

Figure 18: Rockwell 15T Hardness curves for AT72; (T5) aged at 200°C and (T6) solution treated at 420°C and aged at 200°C ...... 56

Figure 19: HAADF-STEM images of AT72 in the (a) as-cast condition and (b) T5 condition (200°C for 65 hours) ...... 57

Figure 20: HAADF-STEM image of AT72 in the T5 condition (200°C for 65 hours) .... 59

Figure 21: Typical tensile curves for each magnesium alloy; AM60B, AT72,AT72 T5(65 hours), and AZ91 ...... 62

Figure 22: Typical tensile curves for AT72 in different treatment conditions ...... 64

Figure 23: Optical micrograph showing the grain size for AT72 in the As-Cast condition and after the T6 heat treatment ...... 65

Figure 24: Predicted solidification path for AT72, ATS0.25, and ATS1.5 using the

Scheil-Gulliver model and PanMg2017 thermodynamic database ...... 72

Figure 25: Backscatter electron image of the as-cast microstructure of gravity cast ATS alloys; AT72, ATS0.5, ATS1.5 ...... 73

Figure 26: Backscatter electron image of the as-cast microstructure of gravity cast AT72

...... 75

Figure 27: Backscatter Electron micrograph of the as-cast microstructure of gravity cast

ATS1.5 ...... 76

Figure 28: Backscatter electron micrograph gravity (a) AT72 and (b) ATS1.5 that were solution treated (420°C for 10 hours) ...... 78

Figure 29: T6 aging curves for gravity cast ATS alloys that were solution treated at

420°C for 10 hours and aged at 200°C ...... 79

Figure 30: T5 aging curves for gravity cast ATS alloys that were aged at 200°C ...... 81

Figure 31: Photograph of the ATS casting produced via HPDC. Courtesy of SJTU ...... 83

Figure 32: Predicted solidification path for ATS and AT72 using the Scheil-Gulliver model and PanMg2017 thermodynamic database ...... 86

Figure 33: Optical micrographs showing the as-cast microstructure of HPDC (a) AT72 and (b) ATS...... 87

Figure 34: Backscatter electron micrograph of HPDC ATS in the as-cast condition ...... 88

Figure 35: HAADF-STEM image of (a) ATS in the as-cast condition and (b) verification of the presence of Mg2Si phase ...... 89

Figure 36: Typical tensile curves for each magnesium alloy; AM60B, AT72, AZ91, and

ATS ...... 93

xvi

Figure 37: Phase fraction vs temperature plot for ATS in both the equilibrium (dotted line) and Scheil (solid line) conditions ...... 95

Figure 38: Backscatter electron micrographs of the as cast and solution treated microstructures of HPDC ATS ...... 100

Figure 39: T6 aging curves for HPDC ATS that was aged at 200°C after solution treated at 420°C for 8 hours and 420°C for 8 hours plus 480°C for 2 hours ...... 102

Figure 40: T5 aging curves for HPDC ATS that was aged at 200°C ...... 105

Figure 41: HAADF-STEM images of ATS in the T5 condition (200°C for 65 hours) .. 106

Figure 42: Energy dispersive x-ray spectroscopy (EDS) maps of Mg, Al, Sn, and Si in

ATS in the T5 condition (200°C for 65 hours) ...... 107

Figure 43: HAADF-STEM of ATS showing orientation of Mg2Sn precipitate ...... 108

Figure 44: Typical tensile curves for AT72 and ATS in the as-cast and T5 condition

(Samples are different geometries and produced with different techniques) ...... 109

Figure 45: Typical tensile curves for ATS in multiple heat treated conditions ...... 111

Figure 46: Optical micrograph showing the grain size for ATS in the As-Cast condition and after a T6 heat treatment ...... 113

Figure 47: Drawing of the fluidity casting ...... 121

Figure 48: Fluidity die inserts cut from P20 steel; (left) ejector side insert (right) cover side insert ...... 121

Figure 49: (a) The 250T Buhler die casting machine (DCM) at OSU. (b) The fluidity die installed in the DCM ...... 123

xvii

Figure 50: Solidification curves for both the Lift380 and EZcast alloys calculated using the Scheil model ...... 126

Figure 51: Conductivity of EZCast and Lift380 calculated using Pandat software ...... 127

Figure 52: Viscosity of EZCast and Lift380 calculated using Pandat software ...... 128

Figure 53: The flow lengths of each wall-thickness for both experimental alloys (melt temperature: 85°C to 105°C superheat, die temperature: 135°C) ...... 130

Figure 54: The cooling rates of the fluidity casting calculated by ProCAST ...... 130

Figure 55: The flow length of each wall-thickness at different fast shot speeds (error bars represent half a standard deviation) ...... 133

Figure 56: Typical Lift380 fluidity castings produced at each shot speed ...... 134

Figure 57: Simulation of fraction solid for both Lift380 (left) and EZCast (right) alloys using the same processing parameters and 577°C metal temperature ...... 136

Figure 58: (Top) Experimental castings produced with Lift380 and (Bottom) simulations of fraction solid based on the process parameters to make the experimental castings ... 137

Figure 59: (Left) Simulation and (Right) experimental casting produced with the sample processing parmeters ...... 138

Figure 60: Drawing of the LIFT Specimen Die casting ...... 146

Figure 61: Top view and side view of the Lift Specimen Die showing the location of the thermocouple pins ...... 147

Figure 62: Drawing of the thermocouple pin inserts, one thermocouple tip is located on the surface while the other is 3mm back form the surface ...... 148

xviii

Figure 63: Lift Specimen Die inserts cut from P20 steel; (Left) ejector side insert (Right) cover side insert with thermocouple pin locations ...... 148

Figure 64: Typically tensile curves for Lift380 tensile specimen produced at different vacuum levels...... 152

Figure 65: Measured cooling curves for each plate thickness ...... 154

Figure 66: Electron backscatter diffraction analysis (EBSD) and grain size distribution of the 2mm plate; alloy: EZCast ...... 156

Figure 67: Electron backscatter diffraction analysis (EBSD) and grain size distribution of the 5mm plate; alloy: EZCast ...... 156

Figure 68: Drawing of side impact bar (SIB) with overall dimensions ...... 158

Figure 69: MAMGASOFT simulation predicting porosity formation in SIB castings for both (Top) EZCast and (Bottom) Lift380 alloys ...... 159

Figure 70: Picture of the SIB casting produced at CAMMET with gating and overflows still attached ...... 160

Figure 71: X-ray results of SIB castings produced with the EZCast alloy ...... 162

Figure 72: X-ray results of SIB castings produced with the Lift380 alloy ...... 162

xix

Chapter 1: Introduction

Light weighting of components is becoming continually more crucial in many industries as companies are prioritizing mass savings to lessen environmental impact, lower costs, and to reduce energy consumption. There are many different techniques that can be used to achieve these weight savings, with the replacing of higher density materials with less dense materials being the most prominent approach. However, this technique has been applied for quite some time now and there are few applications left where a heavier material can be replaced by a currently available less dense material and yet still meet the mechanical property requirements. New material development is essential to further exploit mass savings by substituting heavy materials with lighter ones that have higher specific strengths. Other techniques such as reducing material in the component with novel designs or by using innovative manufacturing processes to manufacture multi-material components, which are able to take advantage of the strength of one material and the density of another, can also be applied to accomplish weight reductions.

Traditionally new alloys were discovered by trial and error methods which was a highly time consuming and costly process. Recent developments have resulted in software which can be used to predict properties of multicomponent alloy systems

1 through thermodynamic models and limited thermodynamic data. This method, called

CALPHAD (CALculation of PHAse Diagrams), can be used to make the development of new alloys more streamlined and cost effective. With faster alloy development and validation, alloys can be created with specific target properties for unique engineering applications.

One method to produce light-weight structural components is the high pressure die casting (HPDC) process. High pressure die casting is a casting process that produces components by forcing molten metal into steel molds, called dies, at very high speeds and pressures. These dies can be manufactured to produce very complex near net shape parts which can be produced reliably at tight tolerances and high volume. These advantages make HPDC a very popular manufacturing process, accounting for a 1/3 of all castings produced. As castings become larger and contain thinner sections, further process development is needed insure that the castings fill and have desirable quality. One such advancement is the introduction of vacuum in the die cavity during the HPDC process, called Super Vacuum Die Casting (SVDC). With the absence of air in the die cavity, the porosity found in castings can be limited and the filling of castings improved.

Development of these processes is needed to insure the success of SVDC and also to provide validation of the mechanical properties of components produced by this method.

This work will pursue the advancement of magnesium alloys for use in high pressure die casting while also providing process development to further the growth of

HPDC for production of structural thin-walled components. Chapter 2 will comprise of background information on the current magnesium alloy systems and the high pressure

2 die casting process. In Chapter 3, the development of an Mg-Al-Sn alloy will be explored, with the goal being to produce an alloy that possesses a better combination of properties compared to current HPDC magnesium alloys. Chapter 4 will describe further advancement of the Mg-Al-Sn alloy system with the addition of Si to improve the mechanical properties and the effectiveness of heat treatments on the alloy. With the advances being made in HPDC industry with state-of-the-art techniques such as SVDC, the heat treatment of alloys is now possible and thus the introduction of heat treatable die casting alloys is essential to capture further increases in properties. Chapter 5 will contain many different topics relating to the advancement of the HPDC process, focusing on

SVDC. The validation of SVDC and how it affects the porosity and mechanical properties is imperative to the growth of this process in industry. Other studies, such as the study of the fluidity of alloys and how simulation software can predict this from properties obtained from the material composition will be covered in this chapter. While the individual chapters will include conclusions on each research topic, Chapter 6 will summarize the overall conclusions of this work and provide future directions of research required to continue the growth of SVDC of structural thin-walled components.

Chapter 2: Background

Magnesium alloys have been extensively used in casting processes because of their properties which are beneficial for the process; however only a few casting alloys have employed for commercial use. This section will review magnesium’s basic properties and the strengthening mechanisms present in typical magnesium alloys. It will give a background on the current magnesium alloys being used in the high pressure die casting application and the limitations of these alloys. The basic processing concepts of high pressure die casting will be introduced as well as the differences in molten metal handling between magnesium and aluminum alloys will be described. Lastly the super vacuum die casting process will be reviewed, explaining how this promotes better castings and how the development of new alloys can leverage these advancements to produce large thin-wall components which have improved mechanical properties.

2.1: Properties of Magnesium

Magnesium is the lightest structural metal, making it an excellent candidate in applications where weight is important such as components in the auto and aerospace industries. Because of recent update on regulations in the United States on emissions and

4 fuel economy in the auto industry, automakers have turned to lighter materials to improve millage. It has been reported that a 10% reduction of weight in an consumer automobile can result in a 4-8% increase in fuel economy [1–3]. With the density of magnesium being only 1.74g/cm3 it is 35% lighter than aluminum (2.7g/cm3) and almost 80% lighter than steel (7.86 g/cm3) making it a favorable choice for designers. When compared to polymers, magnesium is only slightly denser while possessing far better mechanical properties. Magnesium is also the 6th most abundant element on earth with resources available worldwide, making it practical for use in high consumption processes at an affordable price [4].

Magnesium has a hexagonal close-packed (hcp) crystal structure with a c/a ratio of 1.624, near that of the ideal c/a ratio of 1.633, Figure 1. Because of this the ductility of magnesium at low temperatures is reduced as there is only one slip system is active, the

{0001} <112̅0> basal system, Figure 2. Pure magnesium experiences basal slip

(0001) < 112̅0 >, [101̅2]twining, and sub-grain formation during the low temperature primary creep regime [5]. This leads to limited formability of magnesium at low temperatures which makes it unfavorable for forming and extrusion processes preformed at room temperatures. At temperatures greater than 250°C, the {101̅0} and {101̅2} pyramidal and {101̅0} prismatic slip planes are activated while twinning is reduced[5–7].

This leads to more formability but because of softening of particular phases in the alloys, creep performance at elevated temperatures is poor.

Figure 1: Schematic of the hcp crystal structure and lattice parameters of magnesium

There are many other properties of magnesium that make it a favorable choice compared to other structural metals. Magnesium can be machined faster than any other engineering material, leading to reduction of machining time and costs [4]. It can also be easily recycled, so machining swarf and cast scraps can be readily reused. Damping is important for reducing vibrations as well as road noise in automobile applications and magnesium possesses a damping capacity which is greater than both steel and aluminum

[2]. One of the most important properties that magnesium possesses is excellent castability. Because of this and magnesium’s limited formability at room temperatures, over 97% of magnesium alloy structural components are currently produced by casting techniques [8]. These properties make magnesium a favorable metal to use, however because of the lack of strength that pure magnesium possesses, it needs to be alloyed in order to be used in most applications. 6

Figure 2: Schematic showing the (a) slip systems and (b) twin systems present in magnesium alloys [9]

2.1.1: Strengthening Mechanisms

There are three main strengthening mechanisms that readily occur in cast magnesium alloys; grain boundary strengthening, solid solution strengthening, and precipitation strengthening [10]. To develop new alloys which possess exceptional mechanical properties, it is important to understand how each mechanism contributes to the strength of the alloy and how to effectively utilize each of them.

Grain boundary strengthening

Grain boundary strengthening is the strengthening that results from the impedance of dislocation motion by grain boundaries. As the grain size decreases, the area of grain

7 boundaries increases and dislocation motion becomes more difficult. The Hall-Petch relationship, Equation 1, describes the grain boundary strengthening effect where d is the average grain size, K is a constant, and 휎0 is the original yield strength.

−ퟏ⁄ퟐ 흈품풃 = 흈ퟎ + 푲풅 Equation 1

The constant K depends on many different things such as temperature, texture, composition, and preparation. This K value is normally higher for hexagonal close- packed (hcp) metals compared to face center cubic (fcc) metals, so grain boundary strengthening is more effective in metals such as magnesium [11].

The grain size can be controlled by the cooling rate of the manufacturing process or by using alloying additions that act as nucleation sites, called grain refiners, to provide a fine and uniform grain structure. In the high pressure die casting process the cooling rates are very high (in the range of 100s °C/s) and thus we expect a significant amount of grain boundary strengthening present in die cast materials in the as-cast condition [12].

Because the cooling rates are so fast, the addition of solute alloying elements does not affect the grain size very much in the as-cast condition of components produced by

HPDC [13]. However in other heat treatment conditions, such as T5 (artificially aged in as-cast condition) or T6 (solution treated + artificially aged), the grain boundary strengthening effect can be reduced because of the growth of grains which occurs when the material is exposed to higher temperatures. This is where the importance of alloying elements comes in for HPDC alloys, as particular solute atoms will suppress the grain

8 growth at varying degrees. Another way that the suppression of grain growth can be achieved is by having thermally stable phases which pin the grain boundaries thus limiting grain growth. Therefore the maximum strengthening contributed by grain boundary strengthening will be found in the as-cast condition because of the fine grain size resulting from the high cooling rates. Alloying elements which suppress grain growth can assist in reducing the loss of effectiveness of this strengthening mechanism in heat treated conditions.

Solid solution strengthening

Solid solution strengthening is a technique where an alloying element is added to a base element to form a solid solution, which can be either interstitial or substitutional.

A substitutional solid solution is the most common type found in magnesium alloys and is when the solute atom replaces a solvent atom in the crystalline lattice. These solute atoms can cause distortions (local stress fields) in the lattice which will impede dislocation motion leading to strengthening of the material. The difference in the atomic radii and moduli between the solute and solvent atoms will affect severity of this distortion. These local stress fields can either attract or repel dislocations and will modify the ease of dislocation motion. The intensity of this effect can be determined by using the size misfit parameter δ = (dc/da)a and modulus misfit parameter η = dG/dc/G where a is the lattice parameter, c is the atomic solute concentration, and G the shear modulus [11].

The strengthening effect can be expressed according to Labusch by [14]:

ퟐ ퟐ ퟐ ퟐ/ퟑ ퟐ/ퟑ 흈풔풔 = 흈풚풔 + 풁푳푮(휹 + 휷 휼 ) 풄 Equation 2

where 휎ys is the yield stress of pure magnesium, 훧L is a constant, and β is between 1/20 and 1/16. This effect will increase with increasing concentrations of the solute atoms until the concentration reaches a point where a second phase is created. Phase diagrams can be used to predict when particular elements are favorable to create a second phase or stay in solid solution. Because most magnesium alloys are not binary alloys, these phase diagrams can be very complex or hard to calculate without the use of the CALPHAD method. These multiple component systems also change the effectiveness of solute atoms to impede dislocation motion as there can be interactions between the different solute atoms in the solid solution.

Precipitation Strengthening

The last primary strengthening mechanism found in cast magnesium alloys is precipitation strengthening. This occurs when a supersaturated solid solution is aged at an increased temperature so that the elements precipitate out of solid solution forming an additional phase. To produce a supersaturated solid solution an element is added at an amount greater than the low temperature solubility limit but at a high temperature so that there is increased solubility of the element. This supersaturated solid solution microstructure is then locked-in by rapidly cooling the material to a low temperature,

10 called quenching. The alloy is then exposed to a moderate temperature to promote the precipitation of these supersaturated solute atoms into fine second phase particles. The fine precipitates that form during the aging process will impede the movement of dislocation throughout the lattice, causing an increase in strength.

The amount strength gain from precipitation strengthening depends on the size, distribution, and volume fraction of precipitates as well as the interface between the precipitate and matrix which can be coherent or in-coherent[11]. The dislocations can interact with the precipitates in two different ways depending on the size of the precipitate. Smaller precipitates can be sheared in half by dislocations, while large incoherent precipitates resist being sheared and instead the dislocation will bow around these larger particles. The bowing of the dislocations is referred to as Orowan strengthening and is described by Equation 3:

퐺푏 휎 ~ Equation 3 푝 휆

where G is the shear modulus of the matrix, b is the burgers vector, and λ is the average interparticle spacing of the precipitates. If the precipitate nucleation rate can be control in the alloy design, the interparticle spacing can be controlled. By increasing the nucleation rate, the number density of the particles will increase which leads to a smaller interparticle spacing and an increase of the strengthening effect. Microalloying (using elements at a very low weight percent) has been an effective method of precipitate

11 strengthening in many aluminum alloys and can be used to strengthen magnesium alloys as well [15–18].

2.1.2: Magnesium Alloys

Aluminum is the most common alloying element for magnesium alloys as it improves upon many properties of magnesium while only having a slightly higher density. Other common magnesium alloying elements include zinc, silicon, manganese, zirconium, and rare earths, however many other elements are also used in limited capacity. The Hume-Rothery rules are typically applied to determine if an alloying element will form a solid solution or if an intermetallic phase will form upon addition.

These rules describe the conditions which need to be met for an elemental addition to dissolve into a parent metal, forming a solid solution. The first rules states that the solute atomic radius must be within 15% of the solvent atom. Magnesium has an atomic diameter of 0.320nm which in near the diameter of many elements, making a large amount of elements favorable to form a solid solution. Further properties such as valency and electronegativity must also be similar or the alloying element will tend to form an intermetallic compound instead of a solid solution. Table 1 lists many elements which are commonly used in magnesium alloying and their solubility in magnesium.

Table 1: Solubility data (for binary Mg-x systems) for common magnesium alloying additions[19]

Solid solubility Element At. % Wt. % System Aluminum 11.8 12.7 Eutectic Calcium 0.8 1.4 Eutectic Cerium 0.1 0.5 Eutectic Lithium 17.0 5.5 Eutectic Manganese 1.0 2.2 Peritectic Silver 3.8 15.0 Eutectic Thorium 0.5 4.8 Eutectic Tin 3.4 14.5 Eutectic Yttrium 3.8 12.5 Eutectic Zirconium 1.0 3.8 Peritectic Zinc 2.4 6.2 Eutectic

Cast Magnesium Alloys

The majority of cast magnesium alloys can be divided in to two major alloy groups, those containing aluminum and those not containing aluminum but instead containing a small amount of zirconium with various other elements (zinc, silver, rare earths). Zirconium is used as a very effective grain refiner for magnesium as it has the same crystal structure and an almost identical lattice parameter, but binary Mg-Zr alloys do not have the strengths required for most commercial uses [19]. However zirconium cannot be alloyed with typical cast magnesium alloying elements, aluminum or manganese, because it will form detrimental compounds [11]. Instead zirconium is 13 normally added to Mg-RE (rare earth) alloys to refine grains while benefiting from the increased mechanical properties provided from the rare earth additions. Though because of the high cost and low availability of rare earth elements, these alloys are unfavorable for high production casting applications so are normally are processed by sand casting or permanent mold processing techniques instead of high pressure die casting [20].

In die casting, the Mg-Al system of alloys account for nearly 90% of magnesium products because of their great castability, moderate cost, and satisfactory mechanical properties at temperatures below 120°C [20,21]. There are two major Mg-Al alloy systems currently being used for high pressure die castings, Mg-Al-Zn (AZ) and Mg-Al-

Mn (AM). The maximum solubility of aluminum in magnesium at its eutectic temperature of 437°C is 12.7 weight percent while it is less than 2 weight percent at room temperature, shown in the phase diagram in Figure 3. This makes Al additions suitable for both solid solution and precipitation strengthening. As the aluminum content in Mg-

Al alloys increases, the mechanical strength, corrosion resistance, and castability will increase, however the ductility and fracture toughness properties will decrease. Zinc is another addition that can be added to further improve the fluidity and strength of the alloys but needs to be limited to less than 2 weight percent, because at higher content levels the zinc can cause hot tearing to occur. The most widely used Mg-Al alloy is AZ91

(Mg-9 wt.% Al-0.7 wt.% Zn-0.2 wt.% Mn) which is generally used to replace aluminum in strength dominated components (covers, cases, brackets, and housings) that are exposed to ambient temperatures for significant weight savings [19,20,22–24]. When first developed in 1914, AZ91 alloys had problems with corrosion resistance, however new

14 versions of the alloy which limit the impurity content (iron, nickel, copper) have since been developed and improved the performance of corrosion resistance[19].

Figure 3: Mg-Al phase diagram calculated by Pandat

AM50 (Mg- 5 wt.% Al – 0.4 wt.% Mn) and AM60 (Mg- 6 wt.% Al – 0.4 wt.%

Mn) alloys are responsible for the second largest percent of die cast magnesium alloys and are used for structural components where crashworthiness is important (instrument panels, steering systems, and radiator supports) [24]. These alloys have improved ductility over AZ91 because of the reduction of aluminum. When magnesium is alloyed with aluminum, a binary eutectic phase Mg17Al12 forms. This phase, referred to the β phase, forms along the grain boundaries and increases in phase fraction with increasing aluminum content, Figure 4. At high amounts of aluminum, this phase will form 15 continuously around the grain boundaries which will cause brittle fracture and a large reduction in ductility, as observed in the AZ91 alloy. Because of the reduced aluminum content in the AM50/60 alloys, this phase will form discontinuously around the grain boundaries so that the ductility from the matrix can still contribute to the ductility of the alloy. The small addition of manganese leads to an improvement in the yield strength and will also promote the formation of intermetallics with the iron impurities that are present in magnesium alloys. The formation of these Mn-Fe intermetallics will lead to an improvement in the corrosion resistance. These alloys (AM50/60) achieve higher ductility (10-15% elongation) and higher impact strength compared to some common aluminum die cast alloys at a reduced weight making them very popular for light weighting applications [25].

Figure 4: Microstructure of HPDC AM50 [26]

While the mechanical properties of Mg-Al based alloys are good at low temperatures, their strength is greatly reduced at temperatures above 125°C. This can be

16 attributed to the softening of the β-Mg17Al12 phase which has an eutectic temperature of only 437°C[7,20,24]. When at higher temperatures these eutectic phases soften and no longer impede grain boundary sliding, which is the main mechanism of creep at low stress moderate temperatures conditions[27]. Some approaches to improve creep properties of Mg-Al alloys include suppressing the formation of the Mg17Al12 phase, pinning the grain boundary sliding, or slowing the diffusion in the magnesium matrix.

Many alloys are being developed with the goal of improved creep performance so that other automotive applications can be unlocked, such as transmission cases, engine blocks, and engine pistons [28]. However, in this research the creep performance will not be a concern and instead developing a high strength/ductility alloy for use at ambient temperature which can also be further improved with T5 or T6 heat treatment will be the focus.

2.2: CALPHAD

In the early days of alloy development, new alloys were often discovered by inefficient and exhausting ‘trial and error’ methods. These methods were very time consuming and costly, as many different combinations of elements needed to be explored to discover an alloy with desirable properties. While further advancements lead to the use of phase diagrams as a road map to materials development, these phase diagrams were determined by costly and meticulous experiments[29]. In addition experimental phase diagrams could only be determined efficiently for binary and limited ternary systems, as

17 systems with more components were impossible to determine as they were far too complex. Because most real world alloy systems involve greater than two alloying elements, these techniques could not be used readily. Many researchers began to explore the possibility of using computers to calculate binary systems in the 1950’s [29]. In the

1980’s multiple phase diagram calculation software packages became available that could compete phase diagrams for binary and complex higher order systems. With the advancement in software and computing continuing, software packages such as

PANDAT[30], ThermoCalc[31], and FactSage[32] now can be used to calculate phase diagrams and other useful information of complex multi-component systems.

This method of predicting phase equilibria based on limited experimental thermo-chemical data and thermodynamic models to calculate the Gibbs free energy of an alloy system of interest is called the CALPHAD (CALculation of PHAse Diagrams) method [33]. This has led to accelerated development of new alloys and a reduction in costly experiments by allowing for quick assessment of phase equilibria in multicomponent systems and predictions of solidification sequences, phase fraction, heat evolution, transition temperatures, et cetera. With this predicted information alloy designers are able to make educated guesses on what compositions to use when designing a new alloy, thus limiting the amount of experiments.

To predict the Gibbs free energy for each different phase in an alloy, multiple models are used which describe the thermodynamic parameters for each. The Gibbs energy per mole for a liquid or substitutional solid solution is predicted by the following equation[34]:

휙 표 휙 푥푠 휙 퐺푚 = ∑푖 푥푖 퐺푖 + 푅푇 ∑푖 푥푖푙푛푥푖 + 훥 퐺푚 Equation 4

where the first terms describes the Gibbs free energy of the components elements in a reference state, the second describes the ideal Gibbers energy of mixing, and the final term describes the excess Gibbs free energy. The excess Gibbs energy can be described by the Redlich-Kister equation which is limited to 3 parameters at constant T and P[35]:

푥푠 휙 푣=2 푣 훥 퐺푚(푥푖, 푇) = 푥푖푥푗 ∑푣=0 푣 휆(푥푖 − 푥푗푗) Equation 5

where Gm is the molar Gibbs energy, R is the universal gas constant, 휆 is the model parameter and xi/xj are the mole fraction of each component i and j. The intermetallic phases of the system that contain more than one sublattice is described by the compound energy formation[35]:

(푖) (푗) (푖) (푖) (푖) (푖) (푖) 푘 (푖) (푖) 푘 퐺푚 = ∑ 푦푝 푦푞 퐺(푝:푞) + 푅푇 ∑ 푓푖푦푝 푙푛푦푝 + ∑ 푦푝 푦푞 푦푟 푋 ∑푘 퐿(푝,푞,푟)(푦푝 − 푦푞 )

Equation 6

where Gm describes the Gibbs energy as a function of the concentrations of the sublattice species. Again the first term is a reference term, the second term describes the ideal

Gibbs energy, and the third expresses the excess Gibbs energy of the sublattice. The y values are the mole fraction of each species on a sublattice, while the fi is the fraction of a 19 specific sublattice within the crystal and L(p,q:r) are model parameters. These equations and descriptions were acquired from a publication from Luo et.al. [35]. By using these equations the phase equilibria of alloy systems can be calculated with computation software.

One thing to consider is that the CALPHAD approach does not consider the processing conditions used when predicting phase equilibria. Most processing conditions under are non-equilibrium conditions, so the CALPHAD approach has to be coupled with kinetic models to consider these non-equilibrium effects. These kinetic models can also be used to predict solutionizing times, along with precipitations of phases. Prediction of the size distribution, number density, and composition and fraction percent of the precipitate particles in multicomponent alloys is important for heat treatment design.

With coupling the kinetic models with the CALPHAD method, the design of new alloy and the development of heat treatment schedules can be determined more efficiently.

2.3: High Pressure Die Casting

High pressure die casting is a casting process in which liquid metal is injected into a steel mold at high speeds and pressures. The first die cast machine was patented in 1907 by Herman Doehler who was using zinc as the casting material[36]. Now many different materials are used in the high die casting process including zinc, tin, aluminum, and magnesium alloys. There have been many advancements throughout the years which have allowed for complex high integrity castings to be produced with this process.

This process is similar to other casting techniques, such as permanent mold casting, with the main different being the high pressures used to fill the mold. In permanent mold casting, gravity or low pressure is used to fill the die. However in high pressure die casting the metal is forced into the die with a piston at very high pressures

(1000-20000psi), which allows for more complex dies and thinner walled components to be filled at high quality. Because of the high cooling rates associated with the quick filling and the heat transfer of the steel dies, the metal solidifies very quickly allowing for very fast cycle times and thus high volume production. Another benefit of HPDC is that the components produced are near-net shape and can meet very tight tolerances at high repeatability even at high production rates. This reduces machining time on surfaces that need to be very precise, while non-critical surfaces can use the as-cast surface because the steel dies produce a sufficient surface finish.

One major issue with HPDC is that the castings produced have higher levels of porosity compared to other processes and because of this porosity heat treatment of the components in not possible. The formation of this porosity is from a few different sources including gas porosity, shrinkage porosity, and flow porosity. Many solutions have evolved over time to attempt to limit the amount of porosity, such as degassing and vacuum die casting.

In general the high pressure die casting cycle is a simple process. The machine has two platens on which the die blocks are mounted, one is stationary (cover die side) and the other is mobile (ejector side). Hydraulic pressure is used to control the ejector side movement and to clamp the dies closed. Before the die is closed, the two halves of

21 the die are cleaned and lubricated to help with the release of the casting from the die.

Once the die is prepared for casting, the die is closed and securely clamped with high pressure to insure the die stays closed and no molten metal escapes during metal injection. After the dies are closed, molten metal can be transferred into the shot chamber and injected into die with a piston which travels at high speeds (1-10m/s) and high pressures (5-140 MPa). The die fills in a short amount of time (~0.1sec) and high pressure is kept on the metal during solidification. The solidification of the metal is very quick and once complete, the die can be opened. Ejector pins are used to push the part from the die face as shrinkage of the metal or soldering can cause the casting to stick to the die. Once removed, the part is collected and die can be prepared for the next cycle.

These cycle times are very fast compared to other casting techniques and are typically between 30 to 90 seconds depending on the size of the casting and machine.

To produce high pressure die castings there are two different types of machines that can be used; horizontal and vertical, which refers to the travel of the piston relative to the ground. The majority of machines are horizontal die cast machines, while vertical machines are less common they are typically used to make symmetrical components.

There are also two major die castings processes which specify how the metal is transported to the die; hot chamber and cold chamber die casting, which will be discussed in the following sections.

2.3.1: Cold Chamber Die Casting

The cold chamber die casting process is used for materials which have a high melting point or high amount of aluminum; typical metals used in this process are aluminum, brass, and magnesium. The term ‘cold chamber’ comes from the relative temperature different between the shot chamber (also referred to as the shot sleeve) and the holding furnace. In this process the metal is melted to the necessary temperature in a holding furnace and then transported via a hand ladle, auto ladle, or pump into the shot sleeve, as seen in Figure 5. Machines used in the cold chamber process are typically larger than the hot chamber machines and typically have injection pressures between 10-

140MPa. Most cold chamber machines are also horizontal machines. The majority of magnesium automotive die cast components are produced by this process including instrument panels, engine blocks, radiator supports, and seat frames [37].

Figure 5: Schematic of the cold chamber die casting process [38]

2.3.2: Hot Chamber Die Casting

Hot chamber die casting is used for materials which have a low melting temperature such as tin, zinc, lead, and magnesium. The low temperature is required so that the pump that is contained within the metal bath does not get damaged. Magnesium is one of the few metals that can be used in both processes as it has a moderate melting temperature and has a low affinity for iron in the steel components. Even though aluminum alloys possess melting temperatures near magnesium alloys, they cannot be used because they would react with the steel and contaminate the alloy with iron impurities. The major different in this process compared to the cold chamber is the

24 addition of a gooseneck in the shot sleeve and the fact that the shot chamber is contained within the molten metal bath, shown in Figure 6.

Figure 6: Schematic of the hot chamber die casting process [38]

The metal flows freely into this shot chamber when the piston is in the up position. A hydraulic cylinder forces the piston down into the metal which will then flow through the gooseneck into the die cavity. The high pressure is held throughout solidification and once solidification is complete, the piston retracts and more metal enters the shot chamber. This setup allows for quicker cycle times when compared to cold chamber casting, however the injection pressure are limited (5-35MPa). This is because components, such as the nozzle and piston, are at a higher operating temperature

25 and thus experience a reduction in mechanical properties. Because of these limitations, the parts produced by this process are smaller in size comparatively. Commonly produced components included steering wheels, airbag housings, and cases for consumer electronics[37]. Another advantage this process has over the cold chamber process is that the metal has reduced air contact during the transfer from the furnace to die. For magnesium alloys this can be an important benefit as when liquid magnesium is exposed to air, detrimental oxides will form and/or burning will occur.

2.3.3: Molten Magnesium Processing

Molten magnesium will react with oxygen to form oxides or burn when exposed to air. This is because magnesium alloys in the liquid phase do not form a continuous oxide skin which would serve to limit further oxidation. There are two methods in protecting magnesium alloys; flux and fluxless processes. The flux process is typically not used, but can be taken advantage of in situations where fluxless methods do not work.

One such situation when the flux process is used is for alloys that have a melting temperature too high for safe operation with the fluxless method. The most widely used method of protecting magnesium alloys is the fluxless process of using SF6 in a mixture of CO2 or N2. This cover gas is continuously flowed into the crucible and because it is denser than air, it will sit on top of the liquid metal. The SF6 reacts with the liquid Mg to form MgF2 which enhances the MgO film and makes it more protective. It is believed that this is because in the presence of SF6 the contact angle between molten magnesium

26 metal and MgO will be reduced, resulting in the wetting of MgO by magnesium metal and thus forming a more continuous film [39]. However because SF6 has a Global

Warming Potential that is almost 24000 times that of CO2 it is not a very environmentally friendly method. This has resulted in a significant amount of research into other fluxless methods of protecting molten magnesium that have the same effectiveness as SF6 [40–

42]. Until another sufficient method is discovered, SF6 will continue to be the most widely used way to protect molten magnesium.

Another issue that has to be considered because of the burning of magnesium in air is the transport of the molten metal to the shot chamber. In the hot casting chamber process, this is not a concern as the furnace can be protected with cover gas and the metal is never exposed to air because of the gooseneck design. However in the cold chamber casting process the metal is exposed to air during typical ladling operations so this method is ineffective for magnesium. Instead of a ladling method, a pump system with a feed tube is commonly used. This works by installing a mechanical pump onto the furnace which has a feed tube that is inserted into the molten magnesium. When activated the pump forces the molten magnesium through the feed tube which leads to the pour hole in the shot sleeve. This allows for the molten magnesium to be transported to the shot sleeve without being exposed to air and thus prevents the burning of the magnesium alloys during transport.

2.3.4: Molten Aluminum Processing

When molten aluminum alloys are exposed to air, a thin continuous oxide layer will form which will provide protection against further oxidation. This allows aluminum alloys to be transported in air, either by hand ladling or robotic ladling techniques, without the risk of burning. The robotic ladling method is the most widely used technique and is preferred as it is safer and provides for more consistent metal pour volumes.

While the burning of molten aluminum alloys during melt processing is not an issue, aluminum alloys experience issues with the formation of gas porosity because of hydrogen pickup. This gas porosity has adverse effects on the mechanical properties of the castings and can limit the possibility of using heat treatments to further increase mechanical properties. This hydrogen content can be introduced during the casting process in many ways including residual moisture that is on material additions or tools and from moisture in the atmosphere. The residual moisture on tools and material additions can be reasonability controlled, however the moisture content in the atmosphere cannot be and varies significantly based on location and weather conditions. Hydrogen solubility is high in aluminum at temperatures used in holding furnaces and thus the majority of hydrogen found in aluminum alloys is from the moisture in the air.

Magnesium alloys also have a high solubility for hydrogen but do not readily absorb hydrogen at furnace temperatures, therefore the hydrogen content in magnesium alloys

28 can be controlled by reducing residual moisture on the material additions and tools

[43,44].

There are a few different degassing methods that can be employed to reduce the hydrogen content in the liquid aluminum. The most common approaches include lance degassing, rotary impeller degassing, and ultrasonic degassing. In the methods that involve a purge gas, either nitrogen or argon is typically used, with nitrogen being preferred because of the reduced cost. These purge gasses will interact with hydrogen or other inclusions in the melt and bring them to the surface where they can be introduced to the atmosphere or skimmed off the surface.

Lance degassing was the first purge gas method popularized and is a simple process of introducing a purge gas into the melt with a lance made out of iron or graphite.

As the gas leaves the lance, large bubbles are formed around the end of the lance and rise to the surface. This method was slow and not very effective as the large bubbles rise very quickly and only form around a small area at the end of the lance. To improve this, a porous disperser can be added to the end of the lance to generate smaller bubbles and thus a higher total surface area and more of a chance to interact with hydrogen gas. Rotary impeller degassing was introduced to further improve upon this method with the addition of a rotary impeller which spreads the bubbles throughout the melt. This method reduces the hydrogen content more completely and quicker than porous lance degassers. A final method of degassing that does not involve a purging gas is ultrasonic degassing. This method involves inserting a transducer into the melt that induces ultrasonic vibrations.

These vibrations will cause cavitation, rapid formation and collapse of gas pockets, to

29 accelerate the hydrogen diffusion to the surface. This method results in an even more efficient degassing than rotary impeller degassing without the cost of replacing the impeller overtime.[36]

To measure the hydrogen content either a real time hydrogen measurement system or the reduced pressure test is used. The measurement systems employ sensors to give a quick assessment of the hydrogen content in the melt however are costly. The reduced pressure test is the most commonly used method which uses a vacuum system to subject a small cast sample to reduced pressure which will exaggerate the porosity.

Afterwards the density can be measured using the Archimede’s method or the cross sectional area is evaluated and compared to a chart to determine if the cast quality meets standards. However this method not only measures the hydrogen content of the metal but also the presence of inclusions so is more of an overall melt quality measurement instead of hydrogen content measurement [36]. In conclusion, these methods are important for die casters to monitor and improve the aluminum melt quality throughout casting campaigns to insure that satisfactory castings are produced.

2.4: Super Vacuum Die Casting

The major drawback with casting produced by the HPDC process is the increased levels of porosity compared to other casting techniques. The presence of this porosity lowers the ductility of the castings making them unsatisfactory for use in high ductility applications, such as crashworthy components. Porosity also eliminates the possibility of

30 using a solution heat treatment to further improve the mechanical properties of the alloy, as when exposed to high temperatures the porosity will expand and cause blistering in the castings which is detrimental to the mechanical properties. The formation of this porosity can be from a few different sources and can be classified into two major types; shrinkage porosity and gas porosity.

As metals transition from the liquid state to the solid state they undergo a large amount of contraction or solidification shrinkage. This will cause shrinkage porosity to occur in areas where there is not enough metal flow to completely fill voids formed form the solidification shrinkage. Thicker areas in castings are more prone to shrinkage porosity as hot spots form and the regions around solidify initially limiting material flow into the hot spot area. To reduce shrinkage porosity, designers can make adjustments to part design, including additions of cooling or heating lines, or by using alloys that are less susceptible to shrinkage porosity. Cooling lines can be added near hot spots to further increase the cooling rate of these areas to match surrounding areas and thus limit shrinkage porosity formation. While alloys that have a smaller solidification range can also be used to reduce shrinkage porosity as alloys that have larger solidification ranges typically are more sensitive to forming shrinkage porosity because of the longer solidification times. By understanding how the formation of shrinkage porosity occurs, multiple techniques can be utilized to reduce the possibility of the formation of shrinkage porosity in a casting.

Gas porosity is the other type of porosity found in castings and accounts for the majority of porosity in castings produced by the HPDC process. Gas porosity can be

31 caused by three main mechanisms including hydrogen content, die lubricants, and air entrapment. The porosity formed from hydrogen content can be limited by controlling the molten melt quality as discussed previously. Residual lubricants on the die or piston can lead to gasses being created when they are burned off by the molten metal. This can be limited by optimizing the amount of lubricant applied to the die and making sure any blowing step completely eliminates excess lubricant. The final type of gas porosity, air entrapment, is formed when air inside the die is entrapped in the metal due to the turbulent flow conditions. The gating and shot profile can be adjusted to minimize the turbulent flow throughout the die cavity, however this may not be sufficient enough to avoid all formation of porosity.

One way to further reduce the possibility of entrapped air occurring in the casting is to evacuate the die cavity of air prior to filling, referred to as vacuum die casing. There are many different vacuum technologies and methods that are employed to achieve this, but all are similar in that they attempt to reduce the atmospheric pressure in the cavity before the metal enters. Depending on the pressure level attained in the die cavity, the technique is either considered vacuum assisted die casting (60-300mbar) or super vacuum die casting (<60mbar). Super vacuum die casting systems take advantage of advanced controls and monitoring systems, along with powerful vacuum systems, to accomplish these reduced pressures[45].

Typically a vacuum system will use one or more vacuum tanks which are held at low pressures and are much larger in volume than the die cavity. This vacuum tank is connected to the die cavity and the connection between the two is opened after the cavity

32 is sealed, which will occur after the piston passes the pour hole, Figure 7. Because the vacuum tanks are much larger in volume compared to the cavity, once the connection is opened between the two the pressure in the cavity will be reduced very quickly to levels near that of the vacuum tanks. To prevent addition air entering the cavity, a high temperature seal is installed on one of the die faces so that when the die is closed the cavity is completely airtight. The vacuum is then turned off just before the cavity is completely filled in order to reduce the risk of metal flowing into the vacuum system.

There are two different methods to control this mechanism, a valve system or the use of a chill block.

Figure 7: Simple diagram showing the vacuum die casting process; (a) vacuum starts after piston has passed the pouring hole (b) and continues pulling vacuum until casting is filled [46]

The valve system works by physically closing the connection between the vacuum tanks and cavity. It can either open and close based on the stroke of the piston or due to the detected metal pressure. These valve systems are more expensive and more prone to failures; however they have larger cross sectional areas so that a better vacuum can be pulled quicker. The chill block system is a less expensive method that works by promoting metal solidification. Chill blocks utilize a thin corrugated geometry which insures the metal solidifies before going into the vacuum system, Figure 8. Because the small cross sectional area is needed to promote high cooling rates, longer times are needed to pull a vacuum compared to valve systems.

Figure 8: Examples of chill blocks used in vacuum die casting [47]

Castings which are produced with by super vacuum die casting have reduced porosity and thus possess better mechanical properties. They can also be subjected to heat treatments to further increase mechanical properties, as blistering is less likely to occur.

Another advantage of using vacuum in die casting is that it will reduce the back pressure in the die cavity so the metal flow is improved. This allows for larger, more complex, and 34 thinner wall parts to be filled compared to conventional high pressure die casting. Further validation and development of this technique is needed so that a wider range of castings can be produced for commercial use.

Chapter 3: Mg-Al-Sn Alloy Development

3.1: Introduction

The addition of tin (Sn) to magnesium has been explored briefly in the past for its potential to increase strength and creep performance of magnesium alloys at elevated temperatures. Sn is a reasonably low-cost alloying element that has a low melting temperature and is known to improve castability [48]. It possesses a high solubility limit in magnesium of 14.5wt.% at the eutectic temperature of 561°C and 0.45 wt.% at room temperature [49]. This makes it a good candidate for use as a precipitation strengthening alloying addition as a strong solid solution can be formed. In the binary Mg-Sn system, the Mg2Sn phase forms which has a cubic crystal structure with a lattice parameter a =

0.6750nm and is a very brittle and hard phase (~119Hv) [50]. The as-cast microstructure of binary Mg-Sn alloys is dendritic and consists of α-Mg matrix and Mg2Sn [51]. Liu et al. performed a baseline Mg-Sn study in which they varied the composition from 1 to 10 wt.% Sn confirming that the Mg2Sn volume fraction and size increases with increasing amounts of Sn. Increasing additions of Sn lead to further refinement of the secondary dendrite arm spacing of primary Mg which resulted in increasing hardness. The ultimate tensile strength was also observed to increase with increasing Sn up until a composition

36 of 5 wt.% Sn. At concentrations greater than 5 wt.%, the Mg2Sn forms as a semi- continuous network along the grain boundaries that reduces the strength and ductility of the alloy [51]. Therefore optimal mechanical properties were found with a 5wt.% addition of Sn, however the resultant strengths and ductility of this Mg-Sn alloy at room temperature were still lacking compared to commercially available magnesium alloys.

Luo et. al. investigated how the addition of Sn would influence the properties of

Mg-Al based alloys produced by the gravity casting process [52]. The Mg-Al-Sn alloys were prepared with varying amounts of Al (5wt.% to 9wt.%) and Sn (1wt.% to 5wt.%) and were cast into a permanent mold. The microstructure was determined to contain 훼-

Mg, β-Mg17Al12, and Mg2Sn phases. The Mg17Al12 and Mg2Sn phases were observed to increase with increasing Al and Sn content and generally this lead to an increase in the yield strength while decreased the ductility of the alloy. This was because the amount of eutectic phases formed increased with the Al and Sn additions, which reduced the growth of 훼-Mg and reduced the average grain size. It was found that in the permanent mold as- cast condition, an alloy of Mg-7wt.% Al-2 wt.% Sn (referred to as AT72) would produce improved strength and ductility compared to the commercial magnesium alloy AZ91.[52]

Based on the previous research on gravity cast Mg-Al-Sn alloys by Luo et. al., the

AT72 alloy was selected for further investigation for use in high pressure die casting of thin-wall components. It is essential to study the microstructure and mechanical properties of the alloy produced in the high pressure die cast condition, as the fast cooling rates have a significant effect on both. If AT72 can improve upon properties compared to the current magnesium die casting alloys, and compete with the mechanical properties of

37 aluminum alloys, additional weight savings can be captured in structural applications.

There are also additional benefits to using magnesium alloys instead of aluminum alloys in the HPDC process such as reduced affinity to iron, lower specific heat, and better fluidity. As a result of these properties, magnesium alloys have shorter cycle times and longer die lives which make using them even more advantageous as this translates to lowers costs. Additionally thinner wall components are able to be filled with magnesium alloys compared to aluminum alloys (1.0-1.5mm for Mg compared to 2.0-2.5mm for Al) which leads to added weight savings as designers are able to have thin walls in areas where strength is not a concern and add ribs or thicker sections to the structural areas to provide additional strength [37].

This chapter will also explore AT72 in multiple heat treated conditions to attempt to further increase the mechanical properties. Previously studied Mg-Al alloys do not possess an exceptional age-hardening response, especially when compared to aluminum alloys[37,53]. This can be mainly attributed the orientation and coarseness of the

Mg17Al12 precipitates as this makes them inefficient in preventing dislocation motion

[54]. As magnesium is an hcp metal, slip predominantly occurs on the close-packed basal planes (0001) and these continuous Mg17Al12 precipitates tend to form as plates parallel to the basal plane allowing dislocations to easily glide between them. This leads to minimal interaction between the precipitates and dislocations and thus little strength gain is realized. If the number, morphology, or orientation of these precipices can be modified, the age-hardening response can be improved resulting in improved strengths. The addition of Sn and the resulting formation of Mg2Sn may act as heterogeneous

38 precipitation sites for Mg17Al12 or modify the orientation or morphology of the precipitates. This could lead to further reduction of dislocation motion and an increase in strength. The addition of binary Mg2Sn intermetallics could also lead to a reduction in grain boundary motion, which would reduce grain growth during aging. Grain growth during aging leads to a reduction of strength in the alloy and if this can be limited more of the strength increase from the precipitation of fine Mg17Al12 and Mg2Sn phases can be captured.

3.2: Sample Preparation and Methods

To evaluate the microstructure and mechanical properties of AT72, samples were sectioned from door panels that had a thin-wall design throughout the component. These magnesium door panels were produced by high pressure die casting at Eontech in China and were provided by General Motors. Figure 9 shows the overall dimensions of the door panels and the general location of sampling. The average thickness is about 2.25mm throughout the casting. These castings were made with three magnesium alloys for direct comparison; AT72, AM60B, and AZ91. Inductively coupled plasma optical emission spectrometry (ICP-OES) was used to confirm the compositions of each alloy which are presented in Table 2.

Table 2: Compositions found with inductively couple plasma optical emission spectrometry (ICP-OES) of each magnesium alloy; AM60B, AZ91, AT72, and ATS

Composition (weight %) Alloy Al Sn Si Mn Zn Mg ATS 7.46 1.30 0.34 0.33 0 bal. AT72 6.51 1.64 0 0.21 0.19 bal. AM60B 5.50 0 0 0.14 0.13 bal. AZ91 8.16 0 0 0.10 0.69 bal.

The samples subjected to heat treatments were encapsulated in glass tubes at reduced pressures to prevent oxidation occurring during treatment at high temperatures

(400°C+). The samples were quenched in a water bath upon completion of the solution treatment of 420°C for 8 hours. Aging of the hardness samples was performed in an oil bath at 200°C for varying amounts of time (1-300 hours). Hardness values were obtained using the Rockwell 15T hardness scale (1/16 inch steel ball, 15kgf load) and an average of 10 measurements was reported.

To evaluate the microstructure of the alloy in the as-cast and heat treated conditions multiple characterization techniques were used; optical microscope (OM,

Olympus GX71), scanning electron microscopy (SEM, Philips XL-30/FEI Apreo), and transmission electron microscopy (TEM, FEI Tecnai F20 S/TEM). Samples were sectioned from the HPDC door panel and mounted in a typical epoxy resin for both OM and SEM inspection. The mounted specimens were grinded with SiC paper until 1200 grit using standard metallographic practices. Polishing was performed with 3휇m and 1휇m diamond suspensions and then 0.05휇m colloidal silica solution was used for the final

40 polish. To evaluate the microstructure via optical microscopy, samples were etched with an acetic picral solution [70mL ethanol, 10mL acetic acid, 4.2g picric acid, 10mL water].

Figure 9: Photograph showing the dimensions of the HPDC door panel castings and location of tensile bar extraction

For TEM investigation, foils were machined from HPDC specimens using a low speed diamond saw. These foils were mechanically ground to a thickness around 60 - 90

μm and then punched into 3mm discs. The discs were electropolished using a Struers

Tenupol-5 twin-jet unit in an electrolyte consisting of 1500 ml methanol, 300 ml 2- butoxy ethanol, 33.48 g magnesium perchlorate, and 15.9 g lithium chloride at -40°C.

Specimens were further cleaned in a Fischione 1010 ion mill system at an accelerating voltage of 4 KeV.

The mechanical properties of the alloys were found from tensile bars cut directly from the castings using electrical discharge machining (EDM, FANUC). These tensile

41 specimens that are in accordance with ASTM B557 (1’’ gage length, 0.25’’ width) were cut from the door panel in flat areas which had an average thickness of about 2.25mm

[55]. The tensile tests were performed on an Instron 5985 with pneumatic grips (2716-

111) at room temperature. A strain rate of 2mm/mm/min was used to obtain the tensile properties of the alloys and a 25mm strain gage was used (Epsilon 3543) to measure strain. To determine the mechanical properties in the T5 condition, the samples were aged in a conventional furnace at 200°C for 65 hours and 140 hours. The tensile bars for the T4 and T6 testing were encapsulated in glass tubes before undergoing a solution treatment of 420°C for 8 hours and quenched in a water bath. Upon quenching the T6 samples were then aged in a conventional furnace for 100 hours at 200°C.

3.3: HPDC AT72 As-Cast Microstructure

Before investigating the microstructure of HPDC AT72, the solidification path and phase equilibria of the magnesium alloy were reviewed using the CALPHAD method. CompuTherm’s Pandat software and its magnesium thermodynamic database

PanMg2017 were used for the predictions reported in this chapter[56]. The stable and metastable equilibria of a multi-component system can be calculated from its thermodynamic description by using computational methods. In addition to calculating phase diagrams, accurate modeling of microsegregation of alloying elements is essential predicting the solidification path and the as-cast microstructure in alloys [31]. The phase diagram information obtained, such as composition of solid and liquid phases at the 42 solidification interface and the partition coefficients, is essential for microsegregation models. Commercially available CALPHAD software include two models for the prediction of solidification paths of alloys, the equilibrium (Lever) and the Scheil solidification models. The Equilibrium model, also referred to as Lever model, assumes infinite diffusion in liquid and solid states which requires long solidification times and is unlikely in most practical metallurgy processes. The Scheil model on the other hand, assumes no diffusion in the solid and complete mixing in the liquid and assumes a local equilibrium at the solid/liquid interface. The equilibrium and Scheil equations are given below:

푘퐶0 퐸푞푢푖푙푖푏푟푖푢푚: 퐶푠 = Equation 7 1−(1−푘)푓푠

푘−1 푆푐ℎ푒푖푙: 퐶푠 = 푘퐶0(1 − 푓푠) Equation 8

where C0 is the initial alloy composition, Cs is the solidifying solid composition, fs is the solid fraction, and k is the partition coefficient [32]. The partition coefficient affects accuracy of solidification simulations and is strongly dependent on the temperature and concentration of alloying elements. However partition coefficients are not available for multi-component systems, so linearized partition coefficients derived from binary systems are used in solidification simulations. This may introduce significant error in calculations. In general a constant partition coefficient is used in modeling of

43 micosegregation and is a reasonable assumption for binary systems. However, this may not be valid for higher order systems due to the fact that tie-lines may change in these higher order systems. Furthermore, equilibrium phase diagrams can provide solute concentrations in liquid and solid at the solidification interface only for slow and moderate cooling rates. The concentration profile of solute at the interface deviates from equilibrium for high cooling rates where solute trapping takes place. This occurs due to the fact that the solidification front propagates at a rate greater than solute diffusion. This can be taken into account by defining the velocity dependent partition coefficient for high cooling rates [30,31].

Figure 10 shows the solidification path determined by the classical Scheil-

Gulliver model. Again this model assumes two things: complete mixing in liquid and no diffusion in solid; and a local equilibrium at the solid-liquid interface. The cooling rate achieved in most casting processes is considerably faster than equilibrium cooling so the solute atoms do not have time to diffuse. Thus equilibrium solidification is not realistic and segregation occurs in castings. However, at the moving interface it is possible to obtain equilibrium and it is a reasonable assumption to have local equilibrium at the interface[57]. High pressure die casting has very fast cooling rates (order of 100°C/s) and high pressures, which causes the velocity of the interface to be very high and thus the

Scheil model is considered a good approximation as it is not possible for solute atoms to diffuse in the solid. Since Scheil assumes no diffusion in the solid, the microsegregation profile in the solidifying matrix can be approximated by the Scheil model. However, this is really only the case for simple alloy systems. It is important to keep this in mind when 44 looking at Scheil simulation results for alloys produced in HPDC condition, as it may not describe the microstructure completely accurate in complex systems but instead gives a good approximation. Research is being conducted to combat this by including a cooling rate term in the Scheil model for HPDC phase fraction and segregation predictions

[58,59].

Figure 10: The predicted solidification path for Mg-7Al-2Sn using the Scheil-Gulliver model and PanMG17 thermodynamic database

The CALPHAD method predicts the presence of two binary phases in the primary

(Mg) matrix, Mg17Al12 and Mg2Sn, for the AT72 alloy. The primary (Mg) is the first to form at the liquidus temperature of 610°C, followed by the formation of eutectic Mg2Sn at 441°C. The second binary phase, β-Mg17Al12, forms at the solidus temperature of

430°C. Error! Reference source not found.Figure 11 shows the HPDC as-cast 45 microstructure of AT72 alloy. This microstructure is similar to gravity cast microstructure reported previously, which shows the intermetallic phases forming along the grain boundaries discontinuously [60]. The HPDC microstructure is much finer than the gravity cast microstructure because the increased cooling rates experienced in the

HPDC process.

Figure 11: Optical micrograph showing the as-cast microstructure of HPDC AT72

Figure 12 shows the backscattered electron (BSE) image of the as-cast microstructure of HPDC AT72. Backscattered images are created by collecting the electrons which are scattered roughly 180° after interacting with the sample. These electrons are capable of producing images which clearly isolate the different phases present in the microstructure. This is because the heavier elements backscatter electrons

46 more strongly than the lighter elements, making heavier elements appear brighter, because of higher intensity, in BSE images.

When looking at the microstructure in the BSE images, four distinct morphologies and shades can be recognized in the AT72 alloy. The large bright phases in the BSE images are Mn rich intermetallics. Mn is added in small amounts (<0.5%) to most magnesium alloys to improve the corrosion resistance, as it controls the effect of iron impurities found in magnesium alloys by forming Mn,Fe intermetallic compounds. When looking at the predicted AT72 solidification sequence which includes a small addition of

Mn, the Mn-rich intermetallics are predicted to form as the primary phase. These intermetallics will grow to relatively large globular particles. The Mg2Sn intermetallics appear as very small and bright particles which are found in the Mg17Al12 phase, which appears as the second darkest phase in the micrograph. The Mg17Al12 phase has a lamellar morphology while Mg2Sn forms as very small particles within this structure. The

Mg2Sn coexist with the Mg17Al12, suggesting that it acts as a heterogeneous nucleation site for Mg17Al12 as the solidification sequence predicts Mg2Sn to form ahead of the β- phase. Energy-dispersive spectroscopy (EDS) techniques were used to verify the composition of the Mn-rich, Mg2Sn, and β-Mg17Al12 phases found in the microstructure.

Figure 12: Backscatter electron image of the as-cast microstructure of HPDC AT72

Another feature of the microstructure which can be studied is the segregation profile of solute atoms in the grains. The concentration profiles of both Al and Sn in primary (Mg) were calculated with Pandat and are shown in Figure 13. These concentration distributions indicate the segregation of both Al and Sn in (Mg) during solidification of this Mg-Al-Sn alloy. According to the calculation, the center of the grain which is first to solidify, should have concentrations of around 2.2 wt.% Al and

0.75wt.% Sn. As the (Mg) grains grow, the alloying elements in the grain become inhomogeneous and thus the concentrations of Al and Sn increase near the grain boundaries. The concentrations of Al and Sn at the end of solidification, which occurs near the grain boundaries, would be about 11 wt.% and 2.5 wt.% respectively, thus predicting the solute atom segregation seen within the grains in Figure 12. Compositional

48 measurements were performed using EDS methods to confirm the segregation across the grains, which is observed in the BSE images, is mostly solute Al as predicted by the

Pandat software.

Figure 13: Predicted concentration profiles of both (a) Al and (b) Sn in primary (Mg) for the AT72 alloy based on the Scheil-Gulliver model

3.4: Heat Treatment Design and Microstructure

To further improve upon the mechanical properties of magnesium alloys, heat treatments can be utilized. Two heat treatments that can be employed with magnesium alloys, whose nomenclature follows that of aluminum heat treatment designations, are the

T5 and T6 heat treatment schedules. The T5 treatment means that the sample was cooled from an elevated temperature during processing and then artificially aged. Artificially aging is the practice of exposing a material to a moderately increased temperature to accelerate the changes which the microstructure would experience over long periods of time. This allows the microstructure to reach a steady state at a reduced period of time and can lead to precipitation of fine phases. The T6 treatment is a treatment where the sample is subjected to a solution heat treatment, quenched, and then finally artificially aged. Solution heat treatment is the process of exposing a material to a higher temperature at which the liquid phases will not form but the constituents in the microstructure enter into solid solution. The sample is then quenched to lock in the solid solution that has formed. Both of these treatments can be used to modify the properties of alloys and are a common practice for magnesium and aluminum alloys.

The T5 treatment has significant energy savings and less risk of distortion in the components because of the absence of a high temperature solution treatment. T5 is also more commonly used for HPDC components as the high temperatures associated with solution treatment in the T6 treatment can cause blistering in the cast components.

Blistering will lead to a reduction in properties instead of an increase. New techniques, such as super vacuum die casting are being developed to limit porosity and allow for heat treatments to be performed on HPDC components without the risk of blistering.

It is important to understand the microstructures of these treatments so that mechanical properties can be optimized. The first step in designing a T6 heat treatment schedule is to determine the solution treatment needed to achieve complete dissolution of the intermetallic phases. Solution treatment temperatures can be determined by looking at a phase fraction vs temperature plot for the alloy in question. The equilibrium phase fractions vs temperature plot for AT72 was calculated with Pandat and shown in Figure

14. It was found that a solution temperature above 324°C and 372°C would promote the dissolution of Mg2Sn and Mg17Al12 respectively into the matrix. A solution temperature should not exceed 490°C to prevent the formation of any fraction percent of the liquid phase during treatment, as this will lead to reduced mechanical properties. A moderate temperature of 420°C was chosen as a solution temperature for this study to attempt to dissolve the secondary phases while limiting grain growth. Too high of a temperature will promote the growth of grains which will lead to reduced strength, while too low of temperature will not result in complete dissolution and thus limit the amount of solute atoms available for precipitation during the artificially aging step.

Figure 14: The predicated equilibrium phase fraction vs temperature curve for Mg- 7Al-2Sn

To determine the time needed to insure complete dissolution, diffusion simulations created with DICTRA were used to predict the kinetics of the reaction

[31,61]. Figure 15 shows the results of the homogenization simulation of both solute (a)

Al and (b) Sn in (Mg) at 420°C for different holding times. A time of 24 hours was used to represent complete homogenization of solute atoms in the microstructure and it can be observed that a time of 8 hours will produce similar results. Shorter times, such as 5 hours or 2 hours, do not exhibit complete homogenization. Figure 16 shows the solution treated ATS microstructure after holding the specimen at 420°C for 8 hours. After solution treatment the majority of the dendritic secondary phases were dissolved into the

(Mg) matrix and equiaxed grains were formed.

Figure 15: DICTRA simulation showing the homogenization of (a) Al and (b) Sn in primary (Mg) at 420°C held for different hold times

Figure 16: Optical micrograph showing the solution treated microstructure of HPDC AT72

There was a large reduction of the Mg17Al12 beta phase and the segregation of Al and Sn around grain boundaries was completely dissolved into the matrix, Figure 17.

However there was still a small amount of secondary phases located near the grain boundaries which were found to be Mg2Sn and the Mn-Rich particles. These phases are both more thermally stable than Mg17Al12 which lead to the difficulty in completely dissolving these phases as the times calculated by DICTRA were for homogenization of solute Al and Sn atoms in the microstructure. Because the intermetallic phases have reduced solubility in comparison, longer times are needed to dissolve these phases in the microstructure when compared to the dissolution of solute Al and Sn atoms into the matrix.

Figure 17: Backscatter electron image of the microstructure of solution treated (420°C for 8 hours) HPDC AT72

The artificial aging (T5) treatment, which is a prominent treatment to use on

HPDC produced components, was studied as well for AT72. The T5 treatment was performed on samples sectioned from the HPDC AT72 thin-walled castings and a typical magnesium aging temperature of 200°C was used to treat these samples for 1 to 300 hours in an oil bath. The aging hardness curve, Figure 18, shows that the response to the

T5 aging treatment was slow and only a slight increase in hardness was found. There was about a 5% increase from the as-cast hardness to max hardness after being aged for 140 hours, which is too long of a treatment time to use in any industrial application. The alloy was not observed to experience over-aging for the times used in this T5 treatment.

Figure 18: Rockwell 15T Hardness curves for AT72; (T5) aged at 200°C and (T6) solution treated at 420°C and aged at 200°C

The microstructure of AT72 in the T5 condition was observed with TEM to study the precipitation that may have occurred during the aging treatment. Figure 19 shows a

HAADF-STEM (high-angle annular dark field imaging) image of the AT72 in the as-cast and T5 condition. In the as-cast microstructure, the Mg17Al12 and Mg2Sn can be seen forming around the cell boundaries, but there are no fine precipitates observed. After an aging treatment of 200°C for 65 hours which is before peak hardness but well after increase of hardness is found; fine precipitates are observed forming around the cell boundaries. It has been reported that the precipitation strengthening in magnesium alloys is attributable to plate-shaped precipitates found on the basal and prismatic planes [62].

Previous reports also found that the Mg2Sn in Mg-Sn alloys forms as plate and lath morphologies, while Mg17Al12 precipitates in AZ91 were found to be lath and rod-like morphologies [63–65]. Both the Mg17Al12 and Mg2Sn precipitates were identified in the 56

AT72 alloy microstructure and were observed in common morphologies, Figure 20.

There were no fine precipitates found in the middle of the cells, as the precipitates were only found near the cell boundaries.

Figure 19: HAADF-STEM images of AT72 in the (a) as-cast condition and (b) T5 condition (200°C for 65 hours)

Previous research reported that in the T6 condition the fine precipitates were found evenly dispersed throughout the microstructure [66]. However the T6 condition has a solutionizing step which leads to a completely homogeneous microstructure. In the T5 condition this is not the case and the as-cast microstructure is highly inhomogeneous. The higher amount of precipitates on the grain boundary could be because of the segregation of both Al and Sn solute atoms found near the boundaries. The grain boundaries may also act as heterogeneous nucleation sites for these fine Mg2Sn and Mg17Al12 precipitates. 57

This increase in density of these fine precipitates can be attributed for the increase in hardness found in the AT72 alloy after the artificial aging treatment.

During aging the Mg17Al12 phase precipitates out in two modes, discontinuous and continuous precipitation. This is reported in other Mg-Al-Sn and Mg-Al-Zn alloys[54,66]. Continuous precipitation consists of relatively large plates on the basal plane and will tend to form later in the aging process. While discontinuous precipitation ceases early in the precipitation process and is the cellular growth of alternation layers of beta phase and magnesium matrix[54]. Discontinuous precipitates are formed behind a moving grain boundary while continuous precipitates will tend to form in the remaining supersaturated regions[67]. The type of precipitation which occurs also can depend on the aging temperature [68]. Further characterization of these precipitates in AT72 is needed to determine their morphologies and orientation relationships. The fraction percent of

Mg2Sn compared to Mg17Al12 also needs to be investigated so that a complete understanding of the increase in mechanical properties can be obtained.

Figure 20: HAADF-STEM image of AT72 in the T5 condition (200°C for 65 hours)

3.5: HPDC AT72 As-Cast, T5, and T6 Mechanical Properties

Figure 18 shows the aging curve for AT72 in both the artificially aged (T5) and the solution treated plus artificially aged (T6) conditions. The artificial aging temperature for both conditions was 200°C and all hardness samples were aged in an oil bath. The T6 sample was encapsulated in a glass tube and solution treated at 420°C for 8 hours followed by a water quench before aging. The curve shows that the aging response is very slow for AT72. The T5 condition does not reach a peak hardness until around 140

59 hours which is too prolonged for use in commercial applications. There is also only about a 5% increase in hardness observed with the T5 heat treatment.

In the T6 condition, the response is greater and peak hardness is reached in a quicker amount of time. When exposed to a solution treatment, the hardness of AT72 is reduced to about 63 on the Rockwell 15-T scale. After artificial aging the hardness rises around 10% in about 100 hours to hardness of 71. This response is quicker than the T5 condition, however is still too long for use in commercial processing. Previous research on Mg-Sn alloys has also reported slow aging response to aging treatments[69]. The effect of having both Al and Sn precipitates did not greatly affect the aging responses as is shown by the hardness results. Both treatments are also still too long for use in commercial applications and further improvement is needed.

The as-cast mechanical properties for AT72 were compared to current magnesium alloys, AM60B and AZ91, by cutting tensile specimens from thin-wall door inner castings. These castings were produced by the high pressure die casting process and allowed for a direct comparison between the alloys as the samples were produced by the same process and in the same conditions. Each of the tensile specimens were cut from the same general location from the castings which had an average sample thickness of about

2.25mm throughout the gage length of the tensile specimen. The tensile samples were tested with the same equipment and testing parameters and strain was measured with a

25mm extensometer.

Table 3: Mechanical properties of AM60B, AZ91, and AT72 from tensile specimen cut from HPDC thin-wall castings

Alloy YS (MPa) UTS (MPa) Elongation (%)

AM60B 131.7 ± 2.6 239.0 ± 3.7 10.5 ± 0.7

AZ91 158.3 ± 0.5 228.8 ± 3.3 4.0 ± 0.4

AT72 148.9 ± 1.5 250.6 ± 8.1 8.9 ± 1.2

Table 3 contains the yield strength (YS), ultimate tensile strength (UTS) and elongation for each of the alloys. The properties of the HPDC AT72 are greater than the previously reported AT72 gravity cast properties [52]. According to the Hall-Petch relationship, Equation 1, as the grain size decreases the strengthening contribution from grain boundary strengthening will increase. The high pressure die cast samples have reduced grain size compared to gravity cast samples due to the high cooling rates from the processing conditions. This is major reason for the HPDC AT72 yield strengths being improved over those obtained from gravity cast specimen.

The AT72 alloy possesses a yield strength (149MPa) between that of AZ91 and

AM60B, with an elongation (8.9%) which is closer in performance to AM60B. This gives the AT72 alloy an advantage in ultimate tensile strength comparatively. This increase in yield strength can be accredited to the increase in Al and Sn in the alloy. As with an increasing number of binary Mg17Al12 forming along the grain boundaries, a higher increase in strength is observed. With AT72 having aluminum composition which is in between the two alloys you would expect the yield strength to be also. The AZ91 alloy has the highest yield strength because of the large fraction percent of Mg17Al12 however 61 this phase forms a continuous network along the grain boundary which significantly reduces the ductility compared AM60, which has a lower fraction percent of Mg17Al12.

AT72 has a fraction percent of Mg17Al12 which does not form a complete continuous network, thus maintains the majority of ductility found in AM60B. With the introduction of Sn and the Mg2Sn phase in AT72, there is a further increase in yield strength observed.

This leads to an alloy which has a balance of YS and elongation compared to

AZ91/AM60 and higher UTS, Figure 21.

Figure 21: Typical tensile curves for each magnesium alloy; AM60B, AT72,AT72 T5(65 hours), and AZ91

After undergoing a T5 treatment of 200°C for 65 hours, the AT72 alloy experiences a slight increase in yield strength but a large reduction in elongation (3.1%). 62

The formation of precipitates during the aging treatment leads to decreased elongation because they shorten the mean free path of dislocation slip before an obstacle is encountered [66]. The fine precipitates also contribute to an increase in strength via precipitation strengthening as dislocation motion becomes more difficult. However the increase in strength observed from these T5 treatments (65 hours and 160 hours) is not great enough to reach strengths comparable to AZ91. AT72 in the T5 condition of 65 hours at 200°C only has a YS of 150MPa while a poor elongation of 3.1%. When the aging treatment is increased to 160 hours a very slightly higher YS of 151MPa is achieved, but further reduction of elongation (2.7%) is observed, Table 4. These changes are due to a higher fraction percent of Mg17Al12 and Mg2Sn precipitates. Even though these treatments result in an increased of mechanical properties, AT72 still has a yield strength lower than that of AZ91.

Table 4: Mechanical properties of AT72 specimen cut from HPDC thin-wall castings in different heat treatment conditions

Solution Aging time Elongation AT72 YS (MPa) UTS (MPa) Treatment (hours) (%) As-Cast 0 0 148.8 ± 2.5 239.2 ± 10.8 7.7 ± 1.3 T4 420°C 8hrs 0 98.5 ± 0.5 209.1 ± 11.4 7.0 ± 0.8 T5 0 65 150.2 ± 1.3 208.9 ± 2.7 3.1 ± 0.2 T5 0 160 151.3 ± 3.4 206.4 ± 10.5 2.7 ± 1.0 T6 420°C 8hrs 96 115.0 ± 0.8 196.6 ± 20.3 4.4 ± 1.9

When investigating the AT72 alloy in the T4 (solution treated only) and T6 conditions, the mechanical properties of the alloy are reduced, Figure 22. The reduction 63 in elongation can be attributed to the high temperatures of the solution treatment which leads to the expansion of porosity in the samples. Because these castings were made with convention HPDC and not a vacuum assisted die casting process, a large amount of entrapped porosity resided in the specimen. Even though blistering was not apparent on the surface of the samples after solution treatment, the reduction of ductility compared to the as-cast condition suggests that it was present within the sample. In the solution treated condition, AT72 experienced a 33% reduction in YS (98.5MPa) due to the dissolution of the strengthening phases Mg17Al12 and Mg2Sn. It was observed in SEM micrographs that the Mg17Al12 experienced compete dissolution, while only partial dissolution of Mg2Sn was realized.

Figure 22: Typical tensile curves for AT72 in different treatment conditions

With the dissolution of these phases, the alloy should have experience an increase in ductility. However because of the expansion of porosity, this was not the case for these

HPDC samples. In the T6 condition, the AT72 alloy regains 17% of the yield strength due to precipitation of fine Mg17Al12 and Mg2Sn during artificial aging. Figure 23 shows the microstructure of AT72 in both the as-cast and T6 condition. Using the average grain intercept method the average grain size was found. In the as-cast condition the alloy had an average grain size of about 9.9μm, while after being exposed to the T6 heat treatment the grain size increased to 12.0μm. This 20% increase in grain size is partly responsible for the reduced yield strength that the alloy possesses in the T6 condition compared to the as-cast condition. The further reduction of the elongation that is observed in this condition is due to a combination of porosity, grain growth, and presence of precipitates.

Figure 23: Optical micrograph showing the grain size for AT72 in the As-Cast condition and after the T6 heat treatment

Heat treatments that included a solution treatment are not suitable for AT72 in the

HPDC condition, as it results in a reduction in mechanical properties. In the T5 condition, the alloy experiences a slight increase in YS and a reduction in elongation leading to a reduced UTS. This leads to AT72 in the T5 condition underperforming compared to as- cast AZ91. However, in the as-cast condition AT72 displays an increased YS compared to AM60, while maintaining the most of the ductility, giving an alloy that has more balanced YS and elongation compared to AZ91 and AM60. Further research into heat treatment schedules is needed to determine a heat treatment schedule that will improved the properties of AT72 casting produced by super vacuum die casting.

3.6: Conclusions

The results on Mg-Al-Sn alloy development have shown the addition of Sn has a positive influence on mechanical properties. The CALPHAD method can be used to predict microstructure composition of new alloys, while also providing guidelines for heat treatment schedules in attempt to further optimize properties. The as-cast microstructure of AT72 produced by the HPDC process consists of Mg, Mg2Sn, and

Mg17Al12 in a dendritic structure. The secondary phases concentrate around the gain boundaries and provide strengthening to the alloy. The as-cast microstructure of HPDC

AT72 is similar to gravity cast microstructure previously reported, while having finer grain structure due to the increased cooling rates found in HPDC when compared to gravity casting. Artificial aging treatments were not very effective on the AT72 alloy and

66 showed slow response times which are not optimal for industrial use. However, TEM techniques were used to see that there was a formation of fine precipitates as a result of the T5 treatment. These fine precipitates, Mg2Sn and Mg17Al12, aggregated around the grain boundaries and were found in lath, plate, and rod like morphologies. After undergoing solution treatment the Mg17Al12 intermetallics were fully dissolved into the matrix, while the Mg2Sn was only partially dissolved because of its higher thermal stability and lower solubility.

The as-cast mechanical properties of AT72 provided an increase of YS compared to AM60B while only having a slightly lower elongation. The high cooling rates of high pressure die casting, lead to increased strength when compared to prior studies of gravity cast AT72, as the grain size was reduced and according to the Hall-Petch relationship this provides additional strengthening. T5 treatments on AT72 lead to a large decrease in ductility, which gave the alloy a lower ductility and strength when compared to AZ91.

More research is should be performed to improve the ageing response of the alloy and gain a further understanding on the precipitate orientation and fraction percent. Micro alloying additions, such as Na and Zn which have showed increased aging response in

Mg-Sn alloys, should studied to determine if they can be used to shorten treatments times and improve response [65,70]. In conclusion, the CALPHAD method can be applied to aid in the development of new magnesium alloys. The AT72 in the as-cast condition shows a good balance in properties and with slight adjustments in the composition, further improvement on properties can be obtained to produce a high strength magnesium alloy capable for use in high pressure die casting of thin-wall structural components.

Chapter 4: Mg-Al-Sn-Si Alloy Development

4.1: Introduction

Following the previous development of Mg-Al-Sn alloys, the influence of silicon

(Si) on the microstructure and mechanical properties of Mg-Al-Sn alloys was investigated. The development of Mg-Al-Sn alloys lead to AT72 which possessed properties between those of popular magnesium alloys AZ91 and AM60B. However, this alloy experienced a very slow response to a T5 aging treatment and still lacked strengths when compared to AZ91. To further increase the mechanical properties of AT72, the microstructure and mechanical properties of Mg-Al-Sn-Si alloys were evaluated in the as-cast, T5, and T6 conditions and will be reported on in this chapter.

Silicon has previously been investigated as an alloying addition in magnesium alloys because of its relatively low cost and high hardness. It is a popular addition when looking into improving creep properties as typically magnesium alloys soften at higher temperatures [71–73]. When Si additions are added to magnesium alloys, an intermetallic compound Mg2Si forms. This intermetallic phase has a high melting temperature

(1085°C), low density, and high hardness[74]. It has a fcc crystal structure and is similar to that of the Mg2Sn. Generally low Si additions in magnesium and magnesium alloys

68 have led to lower ductility and strengths due to the formation of the brittle script like

Mg2Si phase [75,76]. This script like morphology is observed in microstructures of components that have been processed by techniques which have low cooling rates. When this phase is refined to a globular morphology, there is a resulting improvement in mechanical properties. There are two common methods to refine these eutectic Mg2Si phases; either a small addition of grain refiners can be added to the alloy or the cooling rate can be modified by modifying the processing technique. Because the cooling rate is very high in the HPDC process, it is believed that the Mg2Si will form as refined particles instead of the brittle script like morphology [72]. In this chapter the addition of Si to Mg-

Al-Sn alloys will be examined in the gravity cast condition to determine an addition of Si suitable for HPDC conditions. After a satisfactory Mg-Al-Sn-Si composition is determined, the microstructure and mechanical properties of this optimized alloy will then be studied in the HPDC condition.

4.2: Gravity Cast Experimental Procedure

To prepare gravity cast Mg-Al-Sn-Si alloys, commercial Mg-7Al-2Sn, pure Sn, pure Mg, and an Al-50wt.% Si master alloy were used. The nominal compositions of the alloys prepared are listed in Table 5. The materials were melted in a graphite crucible under a protective gas atmosphere of CO2 + 0.5% SF6 in an induction furnace [MTI –

SP25TC]. The melt was stirred several times to ensure complete homogeneity of the alloying elements. After being held at 750°C for 15 minutes, the molten metal was cast

69 into a steel permanent mold that was preheated to a temperature of 200°C. This mold was a simple cylindrical geometry that produced cylindrical ingots that had a length of 90mm and a 19mm diameter.

Table 5: Nominal compositions (wt.%) of the gravity cast Mg-Al-Sn-Si alloys

Alloy Al Sn Si Mn Mg Designation

AT72 7.15 2.00 0.00 0.43 bal. ATS0.25 7.34 1.99 0.25 0.42 bal. ATS0.5 7.28 1.92 0.48 0.41 bal. ATS0.75 7.31 2.01 0.71 0.40 bal. ATS1.0 7.29 1.81 0.92 0.39 bal. ATS1.5 7.39 1.96 1.32 0.37 bal.

These gravity cast alloys were studied in the as-cast, T5, and T6 conditions. For studying the alloy in the T6 condition, the as-cast alloys were sectioned and encapsulated in glass tubes at reduced pressures followed by a solution treatment at 420°C for 10 hours. After solution treatment the cylindrical ingots were quenched in a water bath and then sectioned into 3mm thick discs using electrical discharge machining (EDM,

FANUC). These discs were then aged in an oil bath at 200°C for varying lengths of time

(1 - 600hours). The T5 specimen were sectioned and aged in the same manner as the T6 samples but without the solution treatment and water bath quench. The hardness tests were performed with the Rockwell 15T hardness scale (1/16 inch steel ball, 15kgf load) and the hardness values reported were obtained from an average of 10 measurements.

The microscopy samples were mounted in epoxy and grinded until 1200 grit with SiC abrasive paper. They were then polished using 0.05휇m colloidal silica as the final step.

The microstructures of the as-cast and solution treated samples were characterized using optical microscopy (OM, Olympus GX71) and scanning electron microscopy (SEM,

Philips XL-30). For observation in the optical microscope, the alloys were etched with an acetic picral solution [70mL ethanol, 10mL acetic acid, 4.2g picric acid, 10mL water].

4.3: Gravity Cast Microstructure and Hardness

Figure 24 shows the solidification path of the Mg-Al-Sn-Si alloys calculated using Pandat software and its Mg thermodynamic database PanMg2017[77]. The calculation is based on the classical Scheil-Gulliver model, which assumes complete mixing in liquid and no diffusion in solid. The previous chapter discussed why the Scheil model is used for prediction of the solidification sequences for alloys produced by casting processes. As shown in Figure 24 and discussed in the previous chapter, the solidification sequence of AT72 is as follows: the primary (Mg) will form at 609oC, followed by the

o formation of eutectic Mg2Sn at 441 C and finally the formation of the β-Mg17Al12 phase at 430°C. The addition of Si will lead to the formation of a new binary phase Mg2Si in the alloy, while no ternary phases are predicted to form. As the Si content increases, at some concentration the primary phase in solidification will become Mg2Si instead of (Mg).

Previous research as shown that this Mg2Si is a very hard, low density phase which has a high melting temperature (1085°C) [74]. In the past gravity cast studies of Mg-Al based

71 alloys, typically low Si additions in magnesium alloys have led to low ductility and strength due to the formation of brittle script like Mg2Si.

Figure 24: Predicted solidification path for AT72, ATS0.25, and ATS1.5 using the Scheil-Gulliver model and PanMg2017 thermodynamic database

The hardness of the alloy in the as-cast condition increases with increasing Si addition because of the increase in the fraction percent of Mg2Si. The Rockwell 15T test was used to investigate how the Si addition affected the hardness and an average of 10 measurements was reported. Table 6 shows the hardness increase from 0wt.% Si to a

1.5wt.% Si addition was about 8%. The hardness of the ATS alloys when compared to

AM50 is much higher, having a 25% to 33% increase in hardness. Figure 25 shows the

72 as-cast microstructure that results from the addition of Si and how increasing Si addition leads to higher formation of script like Mg2Si intermetallics in the gravity cast condition.

At the highest addition amount (1.5wt.% Si), the formation of large script like Mg2Si intermetallics is very obvious. Particles in this morphology can lead to a large reduction in ductility in the alloy.

Table 6: As-cast hardness values (Rockwell 15T) for gravity cast ATS alloys and AM50

As-Cast Hardness (15T) AT72 65.2 ± 0.85 ATS0.25 66.3 ± 0.64 ATS0.5 65.9 ± 0.59 ATS0.75 66.4 ± 0.86 ATS1.0 69.3 ± 0.89 ATS1.5 70.0 ± 0.79 AM50 52.8 ± 1.46

Figure 25: Backscatter electron image of the as-cast microstructure of gravity cast ATS alloys; AT72, ATS0.5, ATS1.5

Figure 26 shows the as-cast microstructure of gravity cast AT72. This microstructure is similar to the HPDC AT72 microstructure, which was reviewed in detail in the last chapter, but has a larger grain size because of the slow cooling rates involved in gravity casting. In the backscatter electron image, you can see how the

Mg2Sn may act as heterogeneous nucleation site for the Mg17Al12 as the Mg2Sn is predicted to form first and is surrounded by β-Mg17Al12 phase. In Figure 27, the backscatter electron micrograph of ATS1.5 is presented. Again the Mg17Al12 and Mg2Sn phases are found near the grain boundaries and the Mg17Al12 phase has a lamellar morphology while Mg2Sn forms as very small particles. There are also small amount Mn- rich intermetallics present in the microstructure, as small additions of Mn are added to improve the corrosion resistance of the alloy. Mg-rich intermetallics form as the primary phase during the solidification process and will have a relatively large particle size.

According to the contrast in the back scatter electron (BSE) image, which is sensitive to the atomic weight of the elements, the segregation of Al and Sn solute atoms in the grains is apparent.

Figure 26: Backscatter electron image of the as-cast microstructure of gravity cast AT72

Figure 27 shows the as-cast microstructure of ATS1.5 and how the addition of Si will lead to the formation of Mg2Si as predicted by the CALPHAD method. The Mg2Si phase forms in a large script like morphology for these gravity cast samples, as cooling rates are slow. The presence of this phase in the alloy is bad for the ductility performance.

After evaluation of the microstructure in the gravity cast condition, the phases observed in the alloys agree with the predictions from the CALPHAD modeling.

Figure 27: Backscatter Electron micrograph of the as-cast microstructure of gravity cast ATS1.5

4.4: Gravity Cast T5 and T6 Microstructure and Hardness

To evaluate how increasing Si additions affected the response to heat treatments, the gravity cast samples were investigated in the T5 (artificial aging) and T6 (solution treatment plus artificial aging) conditions. For the gravity cast samples, the solution treatment temperature of 420oC was selected based on the past AT72 studies while a longer time of 10 hours was used to promote further homogenization of the alloys. At this moderate temperature both the Mg17Al12 and Mg2Sn phases could be at least partially dissolved into the (Mg) matrix. In addition, the grain growth of the primary (Mg) can be

76 controlled by using a moderate temperature rather than a higher temperature closer to the formation of the liquid phase. The thermal stability of the Mg2Si phase is very high and there is almost no solubility of Si in the Mg, so it is unlikely that the Mg2Si particles will dissolve into the matrix at these solution temperatures and times.

Figure 28 shows the microstructures of AT72 and ATS1.5 gravity cast alloys after being subjected to a solution treatment at 420°C for 10 hours. When comparing these microstructures with the as-cast microstructures it is obvious that the Sn and Al solute atom segregation near the (Mg) grain boundary area disappears and that the Mg17Al12 and

Mg2Sn undergo partial or full dissolution. EDS analysis showed that the concentration of

Al and Sn in the grains is rather uniform, which support that the solution treatment at

420oC for 10 hours leads to a homogeneous distribution of these solute atoms. The

Mg17Al12 phase at the grain boundary has been completely dissolved into the (Mg) matrix and the small Mg2Sn particles have experience partial dissolution. As expected, the experimental results indicate that it is easier to dissolve the Mg17Al12 phase compared to the Mg2Sn phase. Both the Mn-rich and Mg2Si phases appear to dissolve very little, if at all, due to their comparatively high stability and melting points. High solution temperatures would be needed to achieve complete dissolution of these phases, but at those temperatures formation of the liquid phase and considerable grain growth are both very likely.

Figure 28: Backscatter electron micrograph gravity (a) AT72 and (b) ATS1.5 that were solution treated (420°C for 10 hours)

A common aging temperature found in the aging of magnesium alloys (200oC) was chosen for this study. According to the calculation, the maximum solubility of Al and Sn at 200oC in the (Mg) matrix are about 3.1 wt.% and 0.3 wt.% respectively. Which suggests that most of the Sn and a large fraction of Al contained in solid solution is predicted to precipitate out of solid solution as Mg2Sn and Mg17Al12. Figure 29 shows the age-hardening curves for the ATS and AT72 alloys aged at 200°C after being solution treated for 10 hours at 420°C. This increase in hardness is due to the precipitation of the

Mg17Al12 and Mg2Sn phases, as there was no precipitation of Mg2Si observed in these

Mg-Al-Sn-Si alloys. As the weight percent of silicon increases in these gravity cast alloys, both the solution treated and peak aged hardness increase. The response to aging slightly decreases with increasing Si content but the time to peak hardness is also observed to decrease slightly. ATS1.5 has the highest peak hardness with a hardness measurement of 77, which is a 13.3% increase from the solution treated hardness of 70. 78

The alloy with the lowest addition of Si, ATS0.25, only has a hardness of 73.7 at the peak aged condition, though experienced an 18% increase over solution treated hardness

(66.3). The gravity cast AT72 alloy also experiences around a 17% increase to a peak hardness (72.2) from 61.6. Table 7 contains the T6 hardness data for all the gravity cast

ATS alloys produced in this study. The time to peak hardness in all these gravity cast alloys is still too delayed at 200°C for use in industrial applications. The addition of silicon did seem to have an effect on the time to peak hardness, as the time slightly decreased with increasing Si.

Figure 29: T6 aging curves for gravity cast ATS alloys that were solution treated at 420°C for 10 hours and aged at 200°C

Table 7: Hardness (Rockwell 15T) values of gravity cast ATS alloys in the as-cast, solution treated, and T6 peak aged condition

Hardness (Rockwell 15T) Solution T6 Peaked Alloy As-Cast Treated Aged AT72 65.2 ± 0.9 61.6 ± 1.3 72.2 ± 1.4 ATS0.25 66.3 ± 0.6 62.1 ± 1.0 73.7 ± 1.0 ATS0.5 65.9 ± 0.6 64.2 ± 0.6 74.2 ± 1.0 ATS0.75 66.4 ± 0.7 65.4 ± 0.7 74.7 ± 1.0 ATS1.0 69.3 ± 0.9 66.1 ± 0.8 75.1 ± 0.6 ATS1.5 70.0 ± 0.8 67.9 ± 1.4 77.0 ± 0.6

Another heat treatment response that was studied with these gravity cast ATS alloys was the artificial aging (T5) treatment. In this treatment, the gravity cast samples were aged at 200°C in an oil bath in their as-cast condition. This lead to a 10% increase in hardness for all samples regardless of silicon composition. Figure 30 shows the aging curves and as was the case with AT72, the time to peak hardness is quite long. The hardness was found to level off around the peak hardness time and no over aging was observed in the treatment times used in this study. In fact, at very long times the hardness of these samples was found to once again increase slightly. The results of the hardness tests at different after treatment at different times are listed in Table 8.

Figure 30: T5 aging curves for gravity cast ATS alloys that were aged at 200°C

Table 8: Hardness (Rockwell 15T) values of gravity cast ATS alloys for as-cast, peak T5, and max T5 hardness

Hardness (Rockwell 15T) T5 Peaked T5 Max Alloy As-Cast Aged Hardness AT72 65.2 ± 0.9 72.7 ± 0.9 72.8 ± 0.9 ATS0.25 66.3 ± 0.6 71.6 ± 0.4 73.7 ± 0.8 ATS0.5 65.9 ± 0.6 71.3 ± 0.8 73.0 ± 0.8 ATS0.75 66.4 ± 0.7 73.1 ± 1.2 73.4 ± 0.9 ATS1.0 69.3 ± 0.9 75.0 ± 0.4 75.2 ± 0.4 ATS1.5 70.0 ± 0.8 76.4 ± 0.5 76.4 ± 0.5

The T5 and T6 heat treatment results for the gravity cast ATS alloys show that the addition of Si does not lead to a significantly better response and only results in a slightly quicker response to heat treatment. However, it does show that the addition of Si does positively affect the hardness of the alloys both in the as-cast condition and heat treated 81 conditions. At high Si additions, the Mg2Si phase formed in very large script like morphologies which can be detrimental to ductility of the alloy. However, the accelerated cooling rates that are found in the high pressure die casting process could result in these

Mg2Si phases to form as smaller globular particles. These smaller particles would be less detrimental to the ductility of the alloy while still providing an increase in hardness and strength. In conclusion, the results of the gravity cast studies shows that an addition of silicon would result in an increase in strength to the AT72 alloy. An addition of

0.75wt.% Si was chosen to be added to AT72 for studies in the high pressure die cast condition. This magnesium alloy will be referred to as ATS (Mg-7wt.% Al-2wt.% Sn-

0.75wt.% Si).

4.5: HPDC ATS Sample Preparation and Methods

After the investigation into Mg-Al-Sn-Si alloys produced by gravity die casting and how the Si addition affected mechanical properties and microstructure, an addition of

0.75wt.% Si was chosen for high pressure die cast studies. This alloy will be referred to as ATS (Mg-7wt.% Al-2wt.% Sn-0.75wt.% Si). The high pressure die cast ATS samples were produced at Shanghai Jiaotong University (SJTU), using a vacuum assisted HPDC system. The vacuum system was able to achieve a vacuum of 250-260mBar in the die cavity as the metal entered the gates. The alloy was prepared with commercially pure Mg and various master alloys. The alloy was held at a temperature between 700-710°C and was injected into the die which was maintained around a temperature of 260°C. Figure 31

82 shows a photograph of the castings produced at SJTU, where the 2.5 mm thick plate in the middle was used for all the hardness analysis and the tensile bar of the left with a 2 inch gage length and 0.25 inch gage diameter was used for mechanical testing (ASTM

B557 Die Cast Specimen)[55]. The chemical composition of the ATS alloy was obtained from inductively coupled plasma mass spectrometry (ICP-OES) and shown in Table 9.

Figure 31: Photograph of the ATS casting produced via HPDC. Courtesy of SJTU

Table 9: Composition of ATS determined by Inductively Coupled Plasma Mass Spectrometry (ICP–OES)

Composition (weight %) Alloy Al Sn Si Mn Zn Mg ATS 7.46 1.30 0.35 0.33 0 bal.

Mechanical properties were obtained by testing the samples in tension with a strain rate of 0.5%/min on an Instron 5985 frame equipped with pneumatic grips (2716-

111). The tests were performed at room temperature and a 50mm strain gage was used

(Epsilon 3543) to measure strain of the specimen. For mechanical properties of the ATS alloy in the T5 condition, the samples were aged in air in a conventional furnace at a temperature of 200°C for 65 hours. The tensile specimen for the T6 conditions were encapsulated in glass tubes to prevent oxidation during solution treatment, quenched in a water bath and then aged in a conventional air furnace at 200°C for 48 hours. Two solution treatment schedules were studied; 420°C for 8 hours and 420°C for 8 hours plus

480°C for 2 hours.

The microstructures of the as-cast and heat treated samples were characterized using optical microscopy (OM, Olympus GX71), scanning electron microscopy (SEM,

Philips XL-30), and Transmission Electron Microscopy (TEM, FEI Tecnai F20 S/TEM).

The samples were sectioned from the HPDC plates and were encapsulated in glass tubes to prevent oxidation during solution treatments. The microscopy samples were subjected to the same heat treatment schedules as the tensile specimens discussed above.

The OM and SEM specimen were mounted in epoxy and polished to a final step of 0.05휇m colloidal silica. OM samples were etched with an acetic picral solution [70mL ethanol, 10mL acetic acid, 4.2g picric acid, 10mL water]. TEM foils were machined from

HPDC specimens using a low speed diamond saw. These foils were mechanically ground to a thickness around 60 - 90 μm and then punched into 3 mm discs. The discs were electropolished using a Struers Tenupol-5 twin-jet unit in an electrolyte consisting of

1500 ml methanol, 300 ml 2-butoxy ethanol, 33.48 g magnesium perchlorate, and 15.9 g lithium chloride at -40°C. Specimens were further cleaned in a Fischione 1010 ion mill system at an accelerating voltage of 4 KeV.

For the aging studies, the hardness samples were cut from the HPDC flat bar castings and aged in an oil bath at 200°C for varying lengths of time (1-300hours). The

T6 specimen were solution treated in glass tubes at the same temperatures and intervals as the tensile specimen. The Rockwell 15T hardness scale (1/16 inch steel ball, 15kgf load) was used for hardness testing and the hardness values reported were obtained from an average of 10 measurements.

4.6: HPDC As-Cast Microstructure and Mechanical Properties

A CALculation of PHase Diagrams (CALPHAD) approach was used to predict the solidification path and phase equilibria of the ATS alloy before microscopy analysis was performed. Pandat software and the PanMg2017 database were used for the calculations [77]. Figure 32 shows the solidification path based on the classical Scheil-

Gulliver model, which assumes complete mixing in liquid and no diffusion in solid. As discussed earlier in the AT72 chapter, this model is not perfect for prediction in the high pressure die cast conditions but can be used to get an estimation of phases present, fraction percent, and transition temperatures.

Figure 32: Predicted solidification path for ATS and AT72 using the Scheil-Gulliver model and PanMg2017 thermodynamic database

The calculated solidification path shows that (Mg) is the primary phase to form at a temperature of 603°C, followed by the formation of eutectic Mg2Si at 591°C. After the formation of the Mg2Si phase, the Mg2Sn and Mg17Al12 phases will form at temperatures of 440°C and 429°C respectively. This is very similar to the solidification path of AT72 where primary (Mg) formation is at 610°C followed by Mg2Sn formation at 441°C and lastly the formation of Mg17Al12 at 430°C. The CALPHAD method predicts the formation of three binary phases and does not predict the formation of any ternary phases.

Figure 33(b) shows the as-cast microstructure of the ATS alloy produced by the high pressure die casting process. This is very similar to the as-cast microstructure of high pressure die cast AT72 alloy also shown in Figure 33(a). The (Mg) matrix is

86 surrounded by Mg17Al12 and Mg2Sn intermetallic phases that form near the grain boundaries, which has also been reported in earlier studies of the AT72 alloys by Shi et. al[66]. The new binary Mg2Si which is found in the ATS alloy is a hard, thermally stable phase. This phase is found to form as particles along the grain boundaries. Earlier in this chapter the microstructure of gravity cast ATS was discussed and micrographs were presented that showed the Mg2Si phase formed in a larger script like morphology. This is because the gravity cast condition has relatively slow cooling rates and thus promotes the formation of particles with large script like morphology. However in the HPDC process, the cooling rate is much higher and the Mg2Si forms predominantly as smaller, globular like particles on the grain boundaries. This change in morphology in the Mg2Sn particles may lead to less of a detrimental effect on the ductility of the alloy. The phases in these micrographs were all identified with scanning electron microscopy and energy-dispersive x-ray spectroscopy techniques.

Figure 33: Optical micrographs showing the as-cast microstructure of HPDC (a) AT72 and (b) ATS

Figure 34 shows the backscatter electron image of the as-cast microstructure of

HPDC ATS. With the heavier elements appearing a brighter shade in BSE images, the aluminum solute atom segregation around the grain boundaries can be seen in the as-cast sample. The secondary phases Mg2Si, Mg17Al12, and Mg2Sn are observed to be located along the grain boundaries. The new binary Mg2Si phase was identified using TEM techniques, shown in Figure 35. The TEM images also confirmed the other intermetallics forming along Mg boundaries and the absence of fine precipitates in the as-cast microstructure. The Mn-rich intermetallics can also be found in the HPDC microstructure, as a small addition of Mn is added to magnesium alloys to improve the corrosion resistance. As discussed in previous chapters, this phase will form as the primary phase and is found as larger particles throughout the microstructure.

Figure 34: Backscatter electron micrograph of HPDC ATS in the as-cast condition

Figure 35: HAADF-STEM image of (a) ATS in the as-cast condition and (b) verification of the presence of Mg2Si phase

Table 10 shows the Rockwell 15T hardness results of HPDC ATS and AT72 alloys in the as-cast condition. The ATS samples were cut from cast flat bars had a thickness of about 2.5mm, while the AT72 samples were cut from the thin-wall door inner casting and had a thickness just under 2mm. While the ATS alloy experienced a slightly slower cooling rate due to the increased thickness, the measured hardness was greater compared to AT72. This is due to the addition of Si and the presence of the very hard globular Mg2Si particles in the alloy. In the gravity cast samples it was observed that as the phase fraction of Mg2Si increased the hardness also increased.

Table 10: Rockwell 15T hardness for HPDC AT72 and ATS in the as-cast condition

Rockwell 15T Hardness Alloy As-Cast AT72 73.5 ± 0.6 ATS 75.7 ± 0.6

The mechanical properties of the ATS alloy are compared to AT72 in Table 11.

The mechanical properties presented in this table were obtained from cast tensile bars which all had a gage diameter of 0.25 inches. However these samples were cast at different locations and the high pressure die casting processing techniques to produce these samples were slightly different. The AT72 tensile bars where produced by two different techniques; both the conventional HPDC and super vacuum die casting (SVDC) processes. While the ATS tensile bars were only produced by one technique, the vacuum assisted die casting process. Vacuum assisted die casting is a vacuum die casting technique which reduces the pressure in the cavity just as the super vacuum die casting technique; however the pressure in the cavity is only reduced to between 60mbar and

300mbar while the SVDC process is capable of reducing the cavity pressure to under

60mbar. The ATS tensile bars which were produced with vacuum assisted die casting for this study had a die cavity pressure of 250-260mbar, while the SVDC ATS had a cavity pressure around 50mbar[66]. When comparing the SVDC AT72 to HPDC AT72, there is a large increase in YS and UTS and a slight increase in elongation. Shi et. al. contributed these increases to the smaller grain size, reduction of porosity, and higher quality of the metal in the SVDC samples compared to the HPDC samples[66]. When comparing the 90 vacuum assisted ATS alloy to SVDC AT72, the properties of the alloys are very similar, with ATS having a slightly lower elongation. While this is not a direct comparison some initial conclusions about the mechanical properties of ATS can be drawn from these results.

If the ATS tensile bars were produced by the SVDC process instead of the vacuum assisted process, one would expect that the YS, UTS, and elongation of the alloy would be improved. This would lead to ATS having an higher YS and UTS compared to

AT72, as the hardness values reported earlier would tend to suggest. This can be attributed to addition of the Mg2Si phase in the ATS alloy, because these hard particles have a small globular like morphology, the strength gain is realized while the alloy only experiences a slight loss in elongation. When comparing the elongation of vacuum assisted ATS to HPDC AT72, the elongations are very similar. Using this relationship, it would be expected that SVDC ATS would have a decreased elongation compared to

SVDC AT72. While the small particle morphology does not lead to an extreme reduction in elongation, they do contribute to a slight reduction. To have a concrete understanding of how the Si addition in ATS improves upon the mechanical properties of AT72, tensile specimen produced with the same processing techniques need to be tested and compared.

Table 11: Mechanical properties of HPDC AT72 and ATS tensile bars

Alloy Process YS (MPa) UTS (MPa) Elongation (%) Vacuum ATS 137.3 ± 3.5 232.2 ±4.8 5.3 ± 0.5 Assisted AT72 [66] HPDC 91.9 175.5 5.3 AT72 [66] SVDC 142.9 231.9 6.4

Figure 36 presents the tensile curve of the ATS alloy compared with the magnesium alloys investigated in Chapter 3. Again these tensile samples were produced with different processing techniques and also have different dimensions. The ATS tensile bars were cast bars with a 6.35mm gage diameter, while the other alloys were cut from thin-wall castings with a thickness around 2.25 mm. When comparing the typical curve for each alloy it can be observed that ATS doesn’t perform that much differently than

AZ91, with having slightly lower yield strength. This could be attributed to the smaller thickness found in the AZ91 samples (~2.25mm) compared to ATS samples (6.3mm diameter) as the reduced grain size would result in higher yield strengths.

Previous studies on how the section thickness affects yield strength of HPDC components shows that as section thickness decreases the YS will increase [78,79]. This can be partially accredited to the reduced grain size due to the higher solidification rates in the thinner sections. The predicted grain structure in the AZ91 (2.25mm thickness) would be smaller than the ATS (6.3mm diameter) grain size. With this in mind, when comparing ATS to AZ91 you could expect that the ATS alloy would possess a larger YS and UTS if the tensile samples of the same geometry were produced with the same processing parameters. The elongation of ATS was observed to be slightly better than

AZ91, however this comparison is hard to make without samples that were produced identically as there are many factors that affect the elongation of HPDC produced specimen.

One last indirect comparison can be made between the ATS and AT72 alloys with these results. It was found that the UTS strength is slightly greater for ATS compared to

AT72, even though the AT72 samples have a smaller grain structure and thus possess a larger increase to YS. This again supports the perception that that ATS would have a greater YS and UTS (plus lower elongation) than AT72 in tensile samples produced with the same technique and dimensions.

Figure 36: Typical tensile curves for each magnesium alloy; AM60B, AT72, AZ91, and ATS

Initial mechanical testing results show that ATS has promising mechanical properties compared to AZ91. When compared in the as-cast conditions, ATS has higher yield and ultimate tensile strengths than AT72, while only a slight reduction in elongation is expected. More studies need to be conducted to investigate the mechanical properties of the alloys with tensile samples that are the same geometry and produced with the same processing techniques. To explore further increases in the mechanical properties of the

ATS alloy, the next section will review the development of three different heat treatment schedules and mechanical properties obtained from these treatments.

4.7: HPDC Heat Treatment Design

The solution treatment response of gravity cast ATS was investigated earlier in this Chapter and it was determined that the 8 hour solution treatment at 420°C did not lead to complete homogenization of the alloy, particularly the dissolution of the Mg2Sn and Mg2Si phases. To facilitate improved dissolution of intermetallics phases in the ATS alloy produced from the HPDC condition, a new solution treatment schedule was designed. To aid in the construction of this heat treatment schedule, a phase fraction vs temperature plot was generated for ATS using the CALPHAD method and both the

Scheil and equilibrium models, Figure 37. Again, the Scheil model is a non-equilibrium model that assumes no diffusion in solid and complete mixing in liquid while also assuming a local equilibrium at the solid–liquid interface. The Scheil model is typically use for processes such as casting, where the cooling rate is quick and does not allow for equilibrium conditions to be met. The equilibrium model assumes complete diffusion in the liquid and solid and is used to describe processes such as furnace cooling which have very slow cooling rates and long times allowing for diffusion to occur.

Figure 37: Phase fraction vs temperature plot for ATS in both the equilibrium (dotted line) and Scheil (solid line) conditions

Looking at Figure 37, the calculated phase fractions from both the Scheil model and equilibrium model can be easily compared. The Scheil model calculations are shown as the solid lines, while the equilibrium model calculation is shown as the dotted lines.

First one can observe that the solidus temperature predicted by the Scheil model (non- equilibrium) is much lower than that by the equilibrium model. This is because in the

Scheil model the microstructure is very inhomogeneous and there are areas with high solute and low solute content. These areas with different solute atom concentrations will have different solidus temperatures, such as in areas with very high Sn solute content the solidus will be lower. In the equilibrium model, the solute distribution will be 95 homogenous and these areas with lower melting temperatures will not exist, leading to a higher overall solidus temperature. It can also be noted in the non-equilibrium condition

(Scheil), the Mg2Sn and Mg17Al12 phases are predicted to be present at temperatures around the solidus temperature, while Mg2Si is a very thermally stable phase and has a much higher melting temperature. In the equilibrium condition the formation temperatures of these phases are much lower temperatures.

During solution treatment, it is important to avoid formation of the liquid phase.

Formation of the liquid phase during heat treatment will result in brittle films to form along grain boundaries or will promote the formation of porosity[80]. Both of these products of grain boundary melting will cause a reduction in mechanical properties of the alloy. High solution temperatures will also result in larger grains as an increase in temperature enables grain growth to occur more readily. The driving force of grain growth is the reduction of internal energy, which is reduced by decreasing the total grain boundary area. The rate at which this occurs is related to many factors including the temperature and time. As the temperature increases, the rate of grain growth will increase as an increase in temperature will increase the diffusion rate. While the longer the sample is held at this increased temperature, the longer enhanced grain growth will occur thus resulting in larger grain size. Therefore the holding times, as well as the treatment temperatures should be limited to lessen the grain growth as a larger grain size will result in a reduced strength. In industry it is also desirable to have shortened heat treatment times so that costs are reduced. With taking all these factors into consideration, the optimization of the holding temperature and time it is very important.

The alloy in this study was produced by the high pressure die casting process, so the Scheil model should be used for phase fraction predictions. The solidus temperature calculated by the Scheil model is 430°C and will be referred to the non-equilibrium solidus temperature (TNEST). It is important that the solution temperature is below this temperature so that the formation of the liquid phase does not occur. However, this temperature is predicted to be too low to promote the dissolution of Mg2Sn and Mg2Si in to the magnesium matrix. To combat this a two stage solution treatment can be taken advantage of, where the first stage is at a lower temperature and the second at a raised temperature. This will promote the dissolution of the lower melting temperature phases first and once they are completely homogenized in to the matrix, the solution temperature can be increased with lesser concerns of the liquid phase forming. The solidus temperature calculated by the lever model (equilibrium) is 500°C and will be referred to as the equilibrium solidus temperature (TEST). This temperature is much higher than the

Mg2Sn phase formation temperature, hence should promote further dissolution of the

Mg2Sn phase into the matrix. The Mg2Si phase, however, has a very high melting point and thus solution treatments under these solidus temperatures are predicted to have little effect on this phase.

A two stage solution treatment of 420°C and 480°C were chosen based on the predicted TEST and TNEST, as well as the predicted Mg2Sn and Mg17Al12 phase formation temperatures. Holding times of 8 hours and 2 hours were selected for these temperatures based on DICTRA simulations used to predict the kinetics of the reaction and the desire to limit time to reduce the amount of grain growth. This solution treatment (420°C for 8

97 hours and 480°C for 2 hours) should allow for complete homogenization of lower melting temperature phases before then exposing the sample to a higher temperature to promote further dissolution of Mg2Sn without the formation of the liquid phase occurring. Further dissolution of these phases, should allow for a better response to the artificial aging treatment and thus increased mechanical properties.

4.8: HPDC Heat Treated Microstructure and Mechanical Properties

Based on the CALPHAD results of the ATS alloy using both the equilibrium and

Scheil models, a two stage heat treatment of 420°C for 8 hours and 480°C for 2 hours was developed to promote further dissolution Mg17Al12 and Mg2Sn. The samples were cut from high pressure die cast bars that had a thickness of 2.5mm. They were encapsulated in glass tubes at reduced pressure for solution treatment and quenched in water after completion of the treatment. Table 12 presents the results of the Rockwell

15T hardness tests that were used to evaluate the hardness of the alloy in both of the solution treated conditions. The two stage treatment reduced the hardness 10% to 67.7 from the as-cast hardness of 74.4, while the one stage only reduced the hardness 6% to

70.3. The two stage solution treatment resulted in almost twice the reduction of hardness compared to the one stage heat. This suggests that further dissolution of the hard intermetallic phase Mg2Sn occurred with the additional step at an increased temperature.

Table 12: Rockwell 15T hardness for HPDC ATS in the as-cast and solution treated conditions

Condition Hardness (15T) % Reduction As Cast 74.4 ± 0.4 - 420 8hrs 70.3 ± 0.5 6% 420 8hrs/480 2hrs 67.7 ± 0.7 10%

The microstructures of both solution treated conditions were observed using scanning electron microscopy, Figure 38. The HPDC as-cast microstructure of ATS consists of Mg17Al12, Mg2Sn, and Mg2Si phases which mostly aggregate near the grain boundaries. After solution treatment at 420°C for 8 hours, the segregated Al and Sn solute atoms in the grains were completely homogenized. The β-Mg17Al12 phase is entirely dissolved into the matrix, while the Mg2Sn phase is only partially dissolved due to the higher melting point and lower diffusivity of Sn. The phases with a high thermal stability, Mg2Si and Mn-rich, experienced no noticeable change in fraction percent after solution treatment. The same trends are observed after exposure to the two stage heat treatment of 420°C for 8 hours and 480°C for 2 hours. The Mg2Sn phase experiences further dissolution, as confirmed with the reduced hardness observed, however is still present in a reduced fraction percent in the microstructure.

Figure 38: Backscatter electron micrographs of the as cast and solution treated microstructures of HPDC ATS 100

After solution treatment, the samples were aged at a typical magnesium aging temperature of 200°C in an oil bath. The aging curves are presented in Figure 39 and show that the two stage solution treatment experiences a greater response to aging. This is because a higher amount of solute Sn and Al atoms dissolved into the Mg matrix and promoted additional precipitation of Mg2Sn and Mg17Al12 during artificial aging. The two stage solution treatment shows a greater response to aging with a 15% increase in hardness from 68.8 to 78.8. The one stage solution treatment only exhibited a 10% increase in hardness from 70.8 to 77.7. Both of the solution treatments experienced peak aging after about 48 hours at 200°C. This is still a quite delayed response and too slow for use in most industrial applications. However the Si addition did promote a faster response to aging, as ATS (48 hours) took about half the time to reach the peak aged condition when compared to AT72 (100 hours). This agrees with the T6 studies that were conducted on the gravity cast ATS samples, as it was shown that increasing Si additions lead to a slight decrease in amount of time to reach the peak aged condition. This faster response to aging observed in the Si containing alloys may be because the Mg2Si intermetallic phase introduces additional heterogeneous nucleation sites for Mg2Sn and

Mg17Al12 precipitates.

101

Figure 39: T6 aging curves for HPDC ATS that was aged at 200°C after solution treated at 420°C for 8 hours and 420°C for 8 hours plus 480°C for 2 hours

Yizhen et al. investigated how the addition of Si to Mg-Al alloys affected the aging response[81]. It was observed that the amount of discontinuous precipitates to form was reduced while the amount of continuous precipitates was increased. This was attributed to the presence of hard, thermally stable Mg2Si particles along the grain boundaries, as these intermetallic phases reduced grain boundary motion and introduced additional dislocations. For discontinuous precipitation, grain boundary motion is required as the grain boundaries serve as a reaction front for precipitation to occur

[82,83]. With the introduction of Mg2Si, the grain boundary motion is limited and thus the amount of discontinuous precipitates is reduced. However continuous precipitation does not required grain boundary motion and instead requires heterogeneous nucleation sites, such as dislocations for precipitation to occur. The introduction of Mg2Si particles

102 was found to produce additional dislocations because the difference in thermal expansion

-6 -1 -6 -1 coefficients between the Mg matrix (29 x 10 K ) and Mg2Si (7.5 x 10 K ) is quite large. When the alloy undergoes temperature changes, the different in coefficients will lead to dislocations forming around the Mg2Si particles to reduce stress. This increased amount of dislocations near the Mg2Si particles act as heterogeneous nucleation sites for continuous Mg17Al12 and Mg2Sn precipitates and leads to an accelerated aging response.

As the Si additions increases, the fraction percent of Mg2Si also increases thus this effect is predicted to become stronger. With the slight addition of Si to the Mg-Al-Sn alloy, we do observe an accelerated response to aging at the treatment temperature of 200°C.

Additionally the increase in hardness found from the aging treatment is also greater for the ATS alloy compared to AT72 because of the additional precipitates.

Another observation in the aging of high pressure die cast samples was that the response was quicker than the response observed in the gravity cast samples. There are a few hypotheses on why this is the case, but they all are based on the premise that precipitation hardening is a diffusion controlled process. Because of the higher solidification rate in HPDC compared to gravity casting, there is a larger amount of vacancies in metals which are produce by HPDC. The faster cooling rate involved leads to a larger concentration of vacancies quenched in during processing. These vacancies can lead to a more efficient aging response as they aid in the diffusion of solute atoms

[84]. Also because of the higher cooling rates, the grain size in HPDC materials is reduced, which also means that there is a higher amount of grain boundary area in the material. The smaller grain size will lead to shorter diffusion lengths which assist in more

103 efficient precipitation. The increase in grain boundary area could lead to further precipitation as there are more grain boundaries available to participate in grain boundary motion promoting discontinuous precipitation. While extra grain boundary area could also lead to the increased formation of heterogeneous nucleation sites for continuous precipitates to form. It is not clear which of these mechanisms is responsible for the quicker response that is observed in the HPDC samples but all may contribute in some fashion.

Figure 40 shows T5 aging curve for HPDC ATS after being aged at 200°C in an oil bath. The alloy only experiences a slight increase in hardness, increasing 4% from the as-cast hardness of 74.4 to the peak aged hardness of 77.7. The peak age condition was found to occur around 65 hours in the ATS alloy. This response was accelerated compared to the T5 aging response of AT72 (140 hours). This is attributed to the addition of Si as discussed earlier in relation to the accelerated T6 response. This increase in hardness is due to fine Mg17Al12 and Mg2Sn precipitates as no fine Mg2Si precipitates were found after any aging conditions of the ATS alloy. The melting point of Mg2Si is much higher than Mg2Sn, so it forms earlier and will grow to large particles located at the grain boundaries. Because of this the amount of solute Si in the matrix is very low and thus the formation of fine Mg2Si precipitates does not occur during aging.

104

Figure 40: T5 aging curves for HPDC ATS that was aged at 200°C

The precipitation of Mg2Sn and Mg17Al12 was observed to occur near the grain boundaries in the T5 condition, Figure 41. The reason for the increased concentration of precipitations near the grain boundary is attributed to the higher amount of solute Al and

Sn near the boundaries from segregation as well as the additional nucleation sites near the grain boundary. Yizhen et al. contributed the additional nucleation sites to the presence of

Mg2Si particles, as the difference in thermal expansion coefficients between Mg2Si and matrix Mg is quite large and results in dislocation formation. These dislocations will act as heterogeneous nucleation sites for Mg17Al12 and Mg2Sn precipitates [81].

105

Figure 41: HAADF-STEM images of ATS in the T5 condition (200°C for 65 hours)

Using EDS techniques the elemental distribution of Mg, Al, Sn, and Si in the T5

ATS microstructure was investigated, Figure 42. The Al rich precipitates were identified as the Mg17Al12 phase, while the Sn rich precipitates were identified as Mg2Sn. Mg17Al12 precipitates were found mostly on the (0001)Mg basal planes and typically had a lath-like morphology. Because of their orientation and number, the blocking of basal dislocation slip in the matrix was rather weak. The Mg2Sn precipitates were found in both of the common orientations, forming as plates parallel to the (0001)Mg basal plane and laths

/rods with the long axis along [2110]Mg, Figure 43. The Mg2Sn precipitates outnumbered the Mg17Al12 precipitates, as shown in the Al/Sn overlap image in Figure 42. The presence of these Mg2Sn and Mg17Al12 precipitates lead to an increase in mechanical properties; however this precipitate strengthening effect was quite weak.

106

Figure 42: Energy dispersive x-ray spectroscopy (EDS) maps of Mg, Al, Sn, and Si in ATS in the T5 condition (200°C for 65 hours)

107

Figure 43: HAADF-STEM of ATS showing orientation of Mg2Sn precipitate

To study the mechanical properties of ATS in the T5 condition, tensile bars were aged at 200°C for 65 hours in a conventional air furnace. Figure 44 shows the typical tensile curves for ATS and AT72 in the as-cast and T5 conditions. These samples are produced by different methods and have different geometries but the response of each alloy to the T5 treatment can be observed. Both alloys experience an increase in yield strength and a reduction in ductility after the T5 treatment. This increase in yield strength was suggested by the hardness results and is due to the formation of Mg2Sn and Mg17Al12 precipitates which make dislocation motion more difficult. The ductility of the alloys is

108 reduced because the presence of additional precipitates shorten the mean free path of dislocation slip before an obstacle is encountered [66].

Figure 44: Typical tensile curves for AT72 and ATS in the as-cast and T5 condition (Samples are different geometries and produced with different techniques)

Figure 45 shows the typically tensile curves for ATS in the T5, T6, and as-cast conditions. Both the two stage (T4-2, T6-2) and one stage (T4, T6) solution treatment schedules were studied. The tensile samples which were subjected to solution treatments were encapsulated in glass tubes at reduced pressures to prevention oxidation from occurring. All the samples were aged in a conventional air furnace at 200°C. Table 13 lists all the different treatment times and temperatures used, as well as the YS, UTS, and elongation for each heat treatment. The T5 treatment leads to about a 5% increase in yield strength for ATS, which is similar to the increase in hardness that was found. However

109 the ultimate tensile strength remains about the same as in the as-cast condition because of a large loss of ductility due to the formation of precipitates.

The solution treated samples should exhibit a large decrease in yield strength, due to the dissolution of Mg17Al12 and Mg2Sn intermetallic phases. The two stage treatment leads to a 38% reduction in yield strength while the one stage treatment only has a 28% reduction. This supports the conclusion that the two stage treatment results in a more complete dissolution of the intermetallic phases as reported by the hardness results.

While the yield strength decreased after solution treatment, the ductility of the alloy did not increase as much as expected. This can be attributed to the expansion of porosity in the samples, as some porosity existed in these samples as they were produced by vacuum assisted high pressure die casting process and not a SVDC process. Blistering was apparent on the surface of some of the two stage tensile bars after solution treatment but not observed on the one stage treatment samples. The ductility of the two stage treatment should be greater because of the reduced amount of intermetallic phases; however it has a much reduced elongation. This further suggests that expansion of porosity occurred in the sample and due to the higher temperatures and longer times of the two stage treatment the blistering was worse and lead to further reduced ductility.

110

Figure 45: Typical tensile curves for ATS in multiple heat treated conditions

Table 13: Mechanical properties of HPDC ATS tensile bars in multiple heat treated conditions

Solution Aging time ATS YS (MPa) UTS (MPa) Elongation (%) Treatment (hours) As-Cast 0 0 137.3 ± 3.5 232.2 ± 4.8 5.3 ± 0.5 T4 420 0 98.7 ± 3.3 247 ± 6.2 8.9 ± 0.7 T4 420/480 0 84.7 ± 4.1 208 ± 17.7 6.5 ± 1.2 T5 0 65 144.0 ± 0.8 230.5 ± 3.8 2.8 ± 0.4 T6 420 48 131.3 ± 2.1 252.7 ± 15.4 6.1 ± 1.3 T6 420/480 48 119.3 ± 6.1 197.3 ± 5.0 2.2 ± 0.3

After aging the solution treated samples the alloy experience increased elongation and yield strengths due to the dissolution of the intermetallic phases and precipitation of 111

Mg2Sn and Mg17Al12 (no detection of Mg2Si precipitates).The yield strength was only slightly reduced compared to the as-cast condition; unlike the AT72 alloy which experienced a very large reduction of YS after T6 treatment. This is because the grain growth in the ATS alloy is reduced compared to the AT72 alloy, which may be due to the presence of the thermally stable Mg2Si particles on the grain boundaries. Figure 46 shows the grain size of ATS in the as-cast and T6 condition. The grain size was measured with the average grain intercept method and was found to be 13.1μm in the as-cast condition and 13.5μm in the T6 condition. This slight increase in grain size (~3%) is much less than the 20% increase in grain size found after the T6 treatment of the AT72 alloy, thus the

ATS alloy experienced a smaller loss in yield strength due to grain growth and retained yield strengths near that of the as-cast condition.

The T6 one stage treatment showed a slight reduction in yield strength while an increased ductility, giving it a higher ultimate tensile strength compared to the as-cast condition. Due to the blistering found in the two stage T6 treatment, after artificial aging the ductility was reduced even more and showed elongations which were worse than the

T5 sample. In conclusion the two stage treatment lead to increased dissolution of the intermetallics phases confirmed by the reduction of yield strength and hardness.

However, because these samples contained a large amount of porosity, the ductility of the alloy was greatly reduced because of the expansion of porosity and blistering during solution treatment. Due to this, the strength gain from the aging treatment was not fully realized and did not lead to much improvement over the as-cast condition. The T5 condition however did display a slight increase in yield strength as well as a large

112 decrease in ductility. If samples are able to be produced by the super vacuum die casting process, thus limiting the formation of porosity, a two stage T6 treatment may be viable to increase properties of the ATS alloy. Additionally further reduction in aging treatment times should be investigated as current times are too long for use in industrial applications.

Figure 46: Optical micrograph showing the grain size for ATS in the As-Cast condition and after a T6 heat treatment

4.9: Conclusions

The investigation of Mg-Al-Sn-Si alloys has shown that the addition of Si has a positive influence on mechanical properties on Mg-Al-Sn alloys. The CALPHAD method can be employed to accurately predict the microstructure composition of new alloys produced by gravity casting and high pressure die casting techniques. CALPHAD was

113 also used to optimize the heat treatment schedules of ATS alloys, leading to further dissolution of intermetallics phases and improvement in mechanical properties. The as- cast microstructure of ATS which was produced by gravity casting consisted of Mg,

Mg17Al12, Mg2Sn, and Mg2Si. These secondary phases aggregated near the gain boundaries and provide strengthening to the alloy. The Mg2Si formed in a large script like morphology, which is detrimental to the ductility of the alloy, because of the slow cooling rates involved with gravity casting. The as-cast microstructure of HPDC ATS is similar to the gravity cast microstructure while having a finer grain structure due to the increased cooling rates. Also in the high pressure die casting condition the Mg2Si particles were refined and thus less detrimental to the ductility, producing an alloy with increased strength with a slight reduction in elongation compared to AT72.

Artificially aging treatments were not very effective on the ATS alloy, however were slightly improved when compared to AT72. The times were still too long for used in industrial applications. Transmission electron microscopy techniques were used to observe the fine precipitates that were a result of the T5 treatment. These precipitates were responsible for the strength gain and identified as Mg17Al12 and Mg2Sn, however no fine Mg2Si precipitates were observed. These fine precipitates aggregated around the grain boundaries and were found in lath, plate, and rod like morphologies in previously reported orientation relationships.

In attempt to improve the dissolution of intermetallics during the solution treatment, a novel two stage treatment was designed with use of the CALPHAD method.

First a temperature below the non-equilibrium solidus temperature (TNEST) was used to

114 dissolve any low melting point phases. After holding for a desirable amount of time as calculated with DICTRA, a second solution temperature just below the equilibrium solidus temperature (TEST) was employed to promote further dissolution of the higher melting point intermetallic phases. After undergoing solution treatment the Mg17Al12 intermetallics were fully dissolved into the matrix, while the Mg2Sn intermetallics were only partially dissolved because their higher thermal stability and lower solubility. The two stage treatment however did show further dissolution of the Mg2Sn phase when compared to the one stage treatment, which was apparent by the further reduction in hardness and yield strength.

The as-cast mechanical properties of ATS provided an increase of YS and UTS when compared to AT72. The ductility of the alloy was reduced slightly with the introduction of Mg2Si intermetallics. After solution treatments, the yield strength was reduced however, the ductility of the alloy did increase because the expansion of porosity in the tensile specimen. This was more apparent in samples that underwent the two stage solution treatment as the higher temperature and longer time lead to blistering on the surface of some tensile samples. After aging, these alloys regained some of the strength due to precipitation strengthening but were still not improved much over the as-cast condition. In the T5 condition, the alloy experienced a slight increase in YS, while maintaining the same UTS because of reduced ductility. The T5 treatment was found to provide the best combination of properties when compared to current magnesium alloys.

Further research is needed for direct comparisons between the AT72 and ATS mechanical properties. Additionally the aging response needs to be investigated more

115 thoroughly so a complete understanding of the orientation and fraction percent of each precipitate can be obtained. Micro alloying additions, such as Na and Zn, should also be investigated to determine if they can be used to shorten treatments times and improve response as they have shown promising results in Mg-Sn alloys [65,70]. In conclusion, the CALPHAD method can be utilized to aid in the development of new magnesium alloys. The ATS alloy in the as-cast condition shows a good balance in properties and has a slight improvement of strength compared to AT72. Unlocking further applications in thin-wall structural components is possible by using the ATS alloy as it shows promising mechanical properties which improve upon the properties of current magnesium alloys.

116

Chapter 5: High Pressure Die Casting Process Development

5.1: Fluidity Studies of Aluminum Alloys

5.1.1: Introduction

Large thin-walled structural castings are becoming increasingly favorable for use in light-weighting applications, as significant mass savings can be achieved when they are produced with high specific strength materials such as magnesium and aluminum alloys. A number of thin-wall castings are currently being produced by the high pressure die casting process and are being used in automotive applications, however more applications are still achievable with further process and alloy development [37,85–87].

When producing these large thin-walled casting, it is essential that the process conditions are optimized as the metal needs to flow further distances from the gate in very thin passages. Material properties are also very important to consider when choosing an alloy for use in these applications as properties such as fluidity and phase transition temperatures are critical for the production of successful castings. The fluidity of an alloy has always been a significant property to consider for use in casting and now as the castings are becoming larger while also thinner it becoming an even more important

117 property. Many studies have been performed on the fluidity of alloys and how it is affected by various casting parameters (melt/mold temperatures, chemical composition, etc.) in the gravity casting process, however few studies exist on these relationships in the high pressure die casting process [88–90]. While there have been limited studies on how the shot speed, wall-thickness, and other parameters relate to the flow distance in high pressure die casting, further investigation is needed [91–93]. In this study a new die with channels of multiple wall-thicknesses was developed to study the effect of key process parameters, such as shot speed, to the fluid flow length of two different aluminum alloys in the high pressure die casting process. For a given alloy this die can also be used to provide optimized process conditions for filling of specific wall thickness sections.

Casting simulation software, ProCAST and MAGMASOFT, were used to design and validate the gating system for the fluidity die and provide initial processing conditions for the die casting trials. Casting simulation is a very valuable tool that can lead to reduced die design times and optimized gating designs as well as provide quicker optimization of processing parameters which produce quality castings. Thermodynamic simulations were also performed using Pandat to predict temperature-dependent alloy properties (specific heat, thermal conductivity, density and viscosity) of the alloy to provide further understanding of the solidification characteristics and fluidity of the alloys being evaluated. With these composition specific properties, the casting software can provide further enhanced predictions for individual alloy systems.

118

5.1.2: Die Design and Experimental Procedure

To study the fluidity of alloys in the high pressure die casting process, a new experimental die was designed and manufactured. Simulation software was used to aid in the design of the die and to insure the gating produced efficient metal flow through the die. Casting simulation software can take a computer-aided design (CAD) model of the casting and with the input of metal properties, mold properties, and processing parameters, is able to predict the flow and solidification of the metal. This can lead to improvements in the quality and yield of the die, while also providing quicker development as the amount of trials to achieve optimized parameters can be reduced.

There are a few different types of casting simulation software that can be used to assist in the design of dies and optimization of processing parameters, though all depend on the same underlying physics. The main differences between the software is how the model is divided and how the numerous material properties and boundary conditions are handled

[94]. Two different simulation software packages were used in the design and study of this die; MAGMASOFT and ProCAST [95,96]. The main difference between these software packages is how the space is divided, as ProCAST uses a Finite Element

Method while MAGMASOFT uses the Finite Volume Method. There are different benefits to each of these methods which are discussed in many publications [94,97].

119

The die designed for these experiments produces a casting that contains three passages which require three-dimensional flow; that is flow which is both parallel (both away from and towards shot sleeve) and normal to the shot sleeve direction, Figure 47.

This design was generated to simulate a more complex part that required metal flow through a three-dimensional geometry, such as a real casting contains, instead of typical fluidity designs for gravity casting which have simpler spiral or step geometries. This die incorporates three passages which each have a different wall thickness, shown as different colors in Figure 47; green - 1mm, blue - 2mm, and yellow - 3mm. These thicknesses where chosen because it covers the range from current thin-wall applications for aluminum alloys (3mm) to future thin-wall applications which could be accessible with further process and alloy development (1mm). Each of these sections has a flow length of 200mm from the gates to the overflow, which was then measured experimentally for each casting produced. This gives a measurable value for fluidity of the alloys which used to studied changes that occur due to modification of processing parameters. Figure 48 shows the die inserts which were cut from P20 steel, for reference of size the large inserts pictured are 7.75 inches by 7.75 inches.

120

Figure 47: Drawing of the fluidity casting

Figure 48: Fluidity die inserts cut from P20 steel; (left) ejector side insert (right) cover side insert

121

Two aluminum alloys were used in these experiments and their compositions that were found by arc spark optical emission spectrometry (SPECTROMAXx) are listed in

Table 14. Pandat software[30] with the PanAl2017 database [56], was used to calculate the solidification curves for the two alloys using the Scheil model. These solidification curves were than validated for both of the alloys by experimental thermal analysis using cooling cup experiments

Table 14: Compositions (wt.%) of the studied alloys obtained by optical emission spectrometry

Alloy Si Fe Cu Mg Mn Ni Zn Sr Ti Al EZCast 8.55 0.11 - 0.28 0.56 0.006 - 0.010 0.008 bal. Lift380 8.83 0.20 3.5 0.29 0.21 0.047 0.11 0.025 0.073 bal.

To produce the fluidity die castings, a 250-ton die casting machine (Bühler H-250

SC) was used. Figure 49 shows the fluidity die installed in the die casting machine which is located in The Center for Design and Manufacturing Excellence (CDME) at The Ohio

State University (OSU). An optimized slow shot speed of 0.38 m/s was used to fill the shot chamber while limiting the amount of entrapped air. The fast shot speed started once the metal reached the gates and was varied from 0.5 m/s to 5 m/s in these experiments.

Previous research on the fluidity of HPDC alloys by Kim et al. reported that the melt quality affects the fluidity of the metal, so each alloy was degassed with a rotary degassing unit (Pyrotek Star 2500) for 20 minutes with argon as the purge gas before

122 casting to insure the cleanliness of the melt [92]. To transport the metal to the shot sleeve a robotic ladling arm (Rimrock) was used with steel ladle that was coated with a boron nitride coating. To experimentally visualize the flow pattern of the metal as the die was filled, partial shots were performed by varying the volume of metal. This was achieved by changing the angle that the ladle exited the molten metal, as a shallower angle leads to a larger metal volume. Once the metal was ladled into the shot sleeve, the shot profile was initiated automatically within a second of pouring. Parameters such as the shot speed, ladle angle, metal temperature, die temperature, and biscuit thickness were documented during each shot for use in comparisons. To measure the flow length of the alloys the total distance traveled from the gate for each different passage was measured and reported on.

Figure 49: (a) The 250T Buhler die casting machine (DCM) at OSU. (b) The fluidity die installed in the DCM

123

5.1.3: Results and Discussion

The compositions of the two alloys that were used in these experiments are presented in Table 14. Using Pandat and the PanAl2017 database, the solidification paths

(temperature vs. solid fraction) for both the EZCast and Lift380 alloys were found and are shown Figure 50. The two alloys are similar in composition with the main difference being a 3.5wt.% addition of Cu, as well as slight additions of Sr and Ti for further grain refinement in the Lift380. The increase in Cu in Lift380 leads to a much lower solidus temperature compared to EZCast, as Cu reduces the eutectic temperature. The slight introduction of Zn will also lead to a further decrease in solidius. The increase in Mn and slight reduction of Si in the EZCast alloy leads to a higher liquidus temperature. The solidus and liquidus of each experimental alloy, Table 15, were found from the predicted solidification paths and were used to select the casting temperature range for each specific alloy.

Lift380 has a very low solidus temperature which leads to the alloy having a large solidification range, more than twice that of the solidification range of EZCast. Large solidification range alloys are more susceptible to form shrinkage porosity. This is because the larger solidification range will result in a larger range that the feeding takes place interdendritically thus promoting formation of microporosity.

124

Table 15: The liquidus, solidus, and fraction of eutectic liquid for each experimental alloy found with Pandat and PanAl2017 thermodynamic database

Liquidus Solidus Solidification Fraction of Alloy (°C) (°C) Range (°C) Eutectic Liquid

EZCast 605 553 52 0.623

Lift380 590 481 109 0.739

The Pandat software was used to predict the simulation input parameters, such as specific heat, thermal conductivity, density, and viscosity as a function of temperature for each alloy. With these properties, the material flow and solidification behavior was able to be more accurately predicted for each specific alloy in the casting simulation software.

Two important material properties to consider when studying the fluidity of an alloy are the viscosity and thermal conductivity. Thermal conductivity describes the materials ability to transfer heat. Materials with higher thermal conductivity will transfer heat quicker than low thermal conductivity materials. Viscosity is the measure of a fluids resistance to gradual deformation via shear or tensile stress; it can also be described as the friction between molecules of a fluid. When a fluid has high viscosity, the material flows slowly and when it has a low viscosity it will flow very quickly, such like water.

Liquid metals have a viscosity that varies greatly with temperature. At temperatures much higher than the liquidus, the liquid metals can flow freely, however at temperatures between the solidus and liquids temperatures the flow slows and it becomes more viscous. Thus it is important to have this material property as a function of temperature for each alloy when looking at flow simulations.

125

Figure 50: Solidification curves for both the Lift380 and EZcast alloys calculated using the Scheil model 126

Figure 51 shows the thermal conductivity of both alloys plotted verse temperature. As a solid, the metal has a very high thermal conductivity but above the solidus when the liquid phase starts to form the conductivity will begin to drop. At the liquidus temperature the conductivity of both alloys is at their lowest. EZCast has a slightly higher conductivity than Lift380, meaning it transfers heat more easily thus will lose heat quicker during solidification in the die.

Figure 51: Conductivity of EZCast and Lift380 calculated using Pandat software

The viscosities of each alloy are plotted verse temperature in Figure 52. Lift380 has a higher viscosity than EZCast at all temperatures, meaning EZCast will flow more easily that Lift380 when at the same temperature. As the temperature increases, the viscosity of both metals decreases. This is an important material property in casting as it describes how easily the liquid metal will flow through the die and because of the fast

127 solidification rates in HPDC this becomes a very important property to consider. Both of these material properties as a function of temperature were used as input parameters in the casting simulation software for a better prediction of filling and flow behavior for each specific alloy.

Figure 52: Viscosity of EZCast and Lift380 calculated using Pandat software

Figure 53 shows the flow length of each thin-wall section for both the EZCast and

Lift380 alloys. These castings were all produced with a fast shot speed of 1.5m/s and a superheat between 85°C and 105°C. Superheat is defined as the difference between the liquidus temperature and temperature at which the metal is poured. Superheat is required in casting so that the alloy does not solidify prematurely in the ladle or shot chamber.

Because of the limited space in the die inserts, oil lines could not be added to the die design. Therefore the die temperature could not be regulated with temperature controlled oil lines and rather was monitored throughout the casting trial with a thermocouple that

128 was inside the die (about 3mm back from casting surface). It was determined that the die inserts were at an average of about 135°C when these fluidity samples were produced for each of the aluminum alloys.

The 3mm section was found to completely fill in the majority of shots at 1.5m/s for both of the alloys. The EZCast alloy experienced a 96% fill rate for the 3mm section, while the Lift380 only experienced a 76% fill rate. In the 1mm and 2mm sections for both alloys, only partial fills were achieved. This is because of the increased cooling rates of these thinner sections. High pressure die casting has very fast cooling rates during solidification and simulations predicted average cooling rates for 1mm, 2mm, and 3mm sections to be in the range of 300°C/sec, 200°C/sec, and 100°C/sec respectively, Figure

54. These cooling rates were calculated as described by Equation 9:

푇 −푇 푐표표푙푖푛푔 푟푎푡푒 = | 푙푖푞푢푖푑푢푠 푠표푙푖푑푢푠| Equation 9 푡푙푖푞푢푖푑푢푠−푡푠표푙푖푑푢푠

where T is the temperature at the liquidus and solidus and t is the time at which the material is at the liquidus and solidus temperatures. Considerable scatter in the flow lengths was found due to the varying casting temperature (85°C-105°C superheat) and die temperature (135 ± 5°C), but by combining all of the shots a few trends in the fluidity of each alloy were found, Table 16. It was observed that the 1mm section averaged about half the flow length compared to the 2mm section for both of the alloys. It seems that the large flow distance difference between the 1mm and 2mm sections is primarily due to the cooling rate difference as these sections were thin enough to limit the flow of metal while 129 the fast cooling rates promoted solidification. In the 3mm section, the combination of a lower cooling rate and larger cross sectional area allowed for the metal to completely fill this passage.

Figure 53: The flow lengths of each wall-thickness for both experimental alloys (melt temperature: 85°C to 105°C superheat, die temperature: 135°C)

Figure 54: The cooling rates of the fluidity casting calculated by ProCAST

130

Table 16: The average flow length for each wall-thickness for both the EZCast and Lift380 alloy at similar superheats and die temperatures Average Flow Length Average Average Die Alloy Superheat Temperature 1mm 2mm 3mm (°C) (°C) EZCast 38.2 71.7 199.5 93.9 141.5 Lift380 52.1 110.8 187.3 95.4 132.3

When comparing the 1mm and 2mm sections, it was observed that Lift380 had a longer average flow length than EZCast for the casting conditions used in these experiments. Because this fluidity casting had a very low shot weight (0.5kg), the fill percent of the shot sleeve was very low (~25% fill). Typically the shot sleeve fill is closer to 50% so that there is a large volume of metal so that the heat loss in the metal in less.

With the small pour volume in these experiments, the superheat that was used was on the low range for making quality castings as there was a lot of heat loss in the shot chamber.

However due to the gas furnace used for these experience, higher temperatures could not be achieved and thus these low superheat temperatures had to be used.

The higher fraction of eutectic liquid and lower solidus temperature of Lift380 compared to EZCast, listed in Table 15, can account for improved flow length in the thinner sections at these low superheat temperatures. A higher fraction of eutectic liquid can help feed the areas which solidify last and generally contributes to increased fluidity in an alloy [36]. The lower solidus temperature allows the metal to flow for more time before solidification is complete. At the lower superheats used in these studies, a lower solidus temperature, thus a larger solidification range, is especially important and will

131 promote greater flow lengths in the thin-wall passages. The 3mm section has a slow enough cooling rate that the casting almost always fills completely for both alloys, therefore it can be said that both alloys show excellent fluidity for the 3mm sections with these parameters. Typical HPDC components have nominal section thickness features interconnected with thin-wall appendages and using the relationship between the section thickness and flow length as shown in these experiments, Table 16 and Figure 53, can be useful to casting design engineers while the detailed casting configuration design phase of new product development.

It should be pointed out that the castability of an alloy is not only dependent on the solidification characteristics, but also the process conditions and die design. In this study, no systematic optimization was done for each alloy. Efforts were made to keep the same superheat (about 95°C above liquidus, max furnace temperature) for both alloys to verify if the current die design could differentiate the relative fluidity of different alloys.

No conclusions should be drawn to the relative castability of the two alloys without additional process optimization of each specific alloy.

Another set of experiments was performed with the Lift380 alloy and consisted of varying the fast shot speed from 0.5m/s to 5m/s while keeping the superheat between 85-

110°C. The flow length for each wall-thickness is plotted against the fast shot speed in

Figure 55. As the shot speed increases from 0.5m/s to 5m/s, the fill time decreases from

75ms to 8ms and the average gate velocities increases from 4m/s to 40m/s. This leads to a roughly linear increase in flow length for each wall thickness besides 3mm which fully fills at each shot speed above 1.5m/s. The response is slightly greater in the 2mm wall

132 thickness compared to the 1mm wall thickness, possibly because of the slightly slower cooling rate found in this thicker passage. Figure 56 shows the typical casting that was produced at each shot speed. Not only does the flow length increase with shot speed, but also the casting quality increases. These results show how shot speed is an important parameter for filling of castings. However, using a very high velocity is not always good for the quality of the casting because the porosity can increased. This is because as the metal velocity increases the metal flow becomes more turbulent and entraps more air.

These faster shot velocities also are harder on the dies and can cause premature failure to the dies which are very expensive. It is important to find a shot velocity that provides a balance between desired flow length that is needed to fill the casting and acceptable casting quality to maintain the components strength.

Figure 55: The flow length of each wall-thickness at different fast shot speeds (error bars represent half a standard deviation)

133

Figure 56: Typical Lift380 fluidity castings produced at each shot speed

Casting simulations were run using ProCAST at the Simulation Innovation and

Modeling Center (SIMCenter) located at The Ohio State University. ProCAST uses a finite element analysis (FEA) code for the modeling of thermal heat transfer (including radiation), fluid flow (including mold filling), and stresses fully coupled with the thermal solution (thermomechanics) [98]. For the simulations reported in this chapter, the following parameters were kept constant unless otherwise noted. A shot profile with a slow shot speed of 0.38m/s and fast shot speed of 1.5m/s was used as used in the experimental trials. Pandat software and databases predicted the thermophysical input

134 properties for both experimental alloys, while temperature dependent HTC values reported by Guo et al. were used, Table 17[99].

Table 17: HTC values used in the ProCAST simulation obtained from Guo et al. [99]

Temperature HTC (°C) (W/m2xK) 0.01 3.0x103 4 Tsol 1.1x10 4 Tliq 2.0x10 1000 2.0x104

For mold thermophysical properties, a similar composition alloy, H13, was selected instead of the P20 that the die was cut in as P20 was not an available die material. Because the heat transfer of these materials are very similar and the stress properties were not consider in these simulation using H13 as a die material is reasonable. Lastly a single cycle simulations were performed because the die temperature was never at a steady state at in the casting trial; instead the recorded die temperature during the selected experimental shots (135°C) was used for a single cycle simulation.

Both alloys exhibited good fluidity and completely filled the die when simulates were run at significantly high casting temperatures. The metal fronts of both alloys during filling at high temperatures were also good. When the casting temperature was lowered to below the liquidus of both alloys, Lift380 was found to completely fill the die while

EZCast had incomplete filling in the 1mm section, Figure 57. In both cases there was early solidification in the 1mm section and there was back flow of metal from the 2mm

135 and 3mm section into the 1mm to ultimately fill the 1mm section in Lift380. The complete filling of the Lift380 casting can again be explained by the lower solidus temperature of the alloy. Because such a low superheat was used in the experiments, better filling was observed in the alloy with a lower solidus and longer solidification range.

Figure 57: Simulation of fraction solid for both Lift380 (left) and EZCast (right) alloys using the same processing parameters and 577°C metal temperature

To observe the metal front as it travels through the die, partial shots were made experimentally with the Lift380 alloy by reducing the metal volume via ladle angle,

Figure 58. As the metal was reduced, the flow length for each section was also reduced.

Experimental processing parameters were recorded during the casting trial and used to create simulations for this case as a comparison. The casting input temperature for the simulation was set at -23°C superheat (567°C) for the Lift380 alloys, taking in account

136 the recorded furnace temperature during casting and the predicted heat loss that was expected before the shot occurred. The simulation results were compared to the experiments by selecting different percent fills that correlated with the experimental partial shots. The flow lengths of the 1mm and 2mm sections were observed to be less than predicted in the simulations, while the 3mm section flow length was greater. Overall the trends of the flow lengths were similar evidenced by increasing wall thickness lead to increasing flow length.

Figure 58: (Top) Experimental castings produced with Lift380 and (Bottom) simulations of fraction solid based on the process parameters to make the experimental castings

137

Another comparison of the metal fill pattern of the casting was made by taking an experimental shot that was conducted at very low temperatures and comparing it to a resulting simulation, Figure 59. To create this casting, a very low furnace temperature and superheat was used (642°C). This lead to premature solidification of the casting, which allowed for the metal front to be observed as filling occurred. When the processing parameters were entered into the casting simulation software, a very similar fill pattern was found when the simulation was at 80% filled. This demonstrates how the software can be used to design gating systems which promote good metal fronts during filling and leads to better overall quality of the castings.

Figure 59: (Left) Simulation and (Right) experimental casting produced with the sample processing parmeters

138

5.1.4: Conclusions

To increase mass savings in lightweight applications, castings are becoming larger and contain thinner walled sections. To fill these large castings, process optimization and alloy fluidity are important factors to study. The fluidity die developed in this study can be used to evaluate fluidity and optimize processing parameters for different alloy systems, by measuring flow lengths for various section thicknesses. The flow length data for each section thickness which was reported in this work can be used as a general guidance in thin-wall casting design. The die shows how the HPDC processing parameters, such as melt temperature, die temperature, and shot speed, can affect the flow length and fluidity of an alloy. Many comparisons can be made between the processing parameters and flow lengths observed in the different wall-thickness passages. Increasing shot speed was shown to increase the flow length during die filling, and this trend was observed to be more significant in thinner wall (1 and 2 mm) than thicker wall (3mm) sections. This is because the increased cooling rates found in thinner sections promoted premature solidification of the alloys. In sections where the cooling rate was not as quick, complete filling was observed for both alloys with the parameters used in this study.

To promote faster and more optimized development of dies, casting simulation software is able to be used to predict metal flow and solidification. This casting simulation software can take a computer aided design (CAD) model of the casting and

139 with input of metal properties, mold properties, and processing parameters is able to predict the flow and solidification of specific metals. This leads to faster development of new die designs and can lead to improvements in the quality of the castings. In these experiments, flow simulations using alloy-specific properties (viscosity, thermophysical data, etc.) for both EZCast and Lift380 alloys where shown to predict the flow length and flow pattern with reasonable accuracy, especially the trend of flow length in the different thin-wall sections. Finally, in the casting conditions used this study (with about the same superheat, and similar die temperature) Lift380 showed longer average flow lengths for 1 and 2 mm sections and similar flow length for 3mm section compared to EZCast. This is likely due to the higher fraction of eutectic liquid and lower liquids in Lift380 as the superheat use in these experiments was quite low. No conclusions should be drawn as to the relative castability of the two alloys without process optimization of each alloy. In conclusion, fluidity of an alloy is a very important material property to consider in the

HPDC process and with the development of this fluidity die, HPDC processing parameters can be studied and accurately compared to simulations.

140

5.2: Super Vacuum Die Casting Development

5.2.1: Introduction

The high pressure die casting process is an economical process which can produce a wide variety of complex near net shape components. Components can be manufactured with high repeatability while being produced at very high production rates because of the high solidification rates involved in the process. The HPDC process has one major disadvantage however, which is the higher probability of porosity formation in the components. The presence of porosity in the castings will lead to reduced mechanical properties, as the porosity will promote crack initiation and propagation[100–103].

Porosity will also reduce the possibility of using heat treatments to further increase the mechanical properties of alloys produced by HPDC. This is because when porosity is exposed to high temperatures, such as those used in heat treatments, it will expand causing blistering in the casting that will further reduce the mechanical properties of the components.

The porosity found in high pressure die castings can be classified into two major types defined by the formation mechanism; shrinkage porosity and gas porosity. When metals transition from the liquid state to a solid state a large amount of contraction, or solidification shrinkage, occurs in the material. This will lead to the formation of shrinkage porosity in areas where there is not enough metal flow to completely fill voids

141 which form due to solidification shrinkage. Shrinkage porosity is normally found in thicker areas of the casting, as these areas are more prone to form hot spots and the regions around these areas will solidify first thus limiting material flow into area the hot spot areas which are the last to solidify. Designers are able to reduce the risk of shrinkage porosity occurring by using simulation software to predict problem areas and making adjustments to the die design such as adding cooling or heating lines to modify solidification rates. Shrinkage porosity can also be limited by using alloys which are less susceptible to solidification shrinkage. Alloys that have large solidification ranges are typically more sensitive to the formation of shrinkage porosity because the solidification takes longer in the problem hot spot areas. Different techniques can be utilized to reduce the probability of shrinkage porosity occurring in a casting; however a small presence of shrinkage porosity in HPDC castings may be unavoidable.

Gas porosity is the main source of porosity in die cast components and can be caused by three main mechanisms; hydrogen content, die lubricants, and air entrapment.

The porosity caused from hydrogen content in the alloy can limited by controlling the molten melt quality as discussed in earlier chapters. Whereas the porosity formed from the burning of residual lubricants on the die or piston can also be limited by optimizing the amount of lubricant applied to the die and making sure the blowing step completely eliminates excess lubricant. By using good practices and optimized process conditions, the porosity formation from the hydrogen content and die lubricants can be reasonably controlled. The final type of gas porosity, air entrapment, is difficult to control as the fast shot speeds involved with HPDC leads to turbulent metal flow conditions. The shot

142 profile and gating of the casting can be optimized to minimize the turbulent metal flow, however because of the intrinsic conditions of the HPDC process it can only be partially limited by modifying these parameters.

One technique that can be applied to further reduce the probability of entrapped air in the casting is to reduce the presence of air in the die cavity before the metal enters the cavity, referred to as vacuum die casting. There are many different technologies and methods that are employed to achieve this, but each attempt to reduced pressure in the die cavity by applying a vacuum [46,104]. Typically a vacuum die casting system will use one or multiple vacuum tanks that are held at very low pressures. These tanks are relatively large in volume compared to the die cavity so when the connection to the die cavity is opened the pressure equalizes very quickly to a low pressure. The die cavity is connected to the vacuum tanks by means of a chill block or vacuum valve on one end and the other end is sealed when the shot piston passes the pour hole. Some systems take advantage of a high temperature polymer seal between the ejector die and cover die to prevent unwanted air from entering the cavity.

Castings produced by the vacuum die casting process have reduced porosity compared to conventional HPDC components. This is because there is a reduction of air in the die cavity that is able to be entrapped by the turbulent metal flow during filling.

Casting that have reduced porosity will possess better mechanical properties and can be subjected to heat treatments as blistering is less likely to occur [26,101,102]. An additional advantage of using vacuum in die casting is that the back pressure in the die cavity is reduced so the metal flow is improved. This allows for larger, more complex,

143 and thinner wall parts to be filled compared with parts manufactured by the conventional high pressure die casting process. Further validation and development of vacuum die casting is needed so that larger thin-walled casting can be produced for commercial use.

This chapter will explore the mechanical properties of aluminum alloys produced by conventional and vacuum die casting techniques. A new die was developed that produces multiple 3 point bend plates and a tensile specimen. The 3 point bend specimen varied in wall thickness from 5mm to 2mm so investigation on how the different thicknesses were affected by the introduced of vacuum could be performed. This chapter will conclude with an investigation into the presence of porosity and how it compared with simulation porosity predictions of an industrial casting that was produced with and without super vacuum die casting.

5.2.2: Die Design and Experimental Procedure

The die fabricated for these studies, which will be referred to as the Lift Specimen die, consists of multiple 3 point bend plates and an ASTM tensile bar, Figure 60. The 3 point bend plate each have a different thickness (5mm, 3mm, and 2mm) so that the microstructure and mechanical properties of different cooling rate can be studied. It also allows for the investigation into how the vacuum level affects the filling and quality of the plates produced at different thicknesses. Simulation software was utilized to optimize the gating design and to insure that the metal front was even throughout the filling of

144 each of the plates. Two simulation software packages were used in the design and study of this die, MAGMASOFT and ProCAST [95,96].

The die set includes a copper chill block, which is a separate insert in the die block, so that the liquid metal solidifies before flowing into the vacuum lines. This chill block is always ‘open’ so the vacuum has to be turned on and off automatically by the system to evacuate the air in the die cavity. The vacuum system turns ‘on’ after the shot is initiated and the piston passes the pour hole, thus creating a closed system. This is controlled with a laser and retroreflector that sends a signal to the vacuum system when the piston is about to pass the pour hole. The vacuum stays on for a programed amount of time that was calculated with the shot profile and simulation software. It is important that the vacuum is turned off before the metal completely fills the die or the vacuum may promote metal flow through the chill block into the vacuum line. This would cause a blockage in the system and would prevent an acceptable vacuum level from being achieved in future shots.

145

Figure 60: Drawing of the LIFT Specimen Die casting

The Lift Specimen die incorporates three different wall thicknesses of 3 point bend bars, which covers from common thicknesses found in aluminum HPDC castings

(5mm) to thin-wall applications (2mm). One unique feature in this die is the ability to insert thermocouple pins to measure the temperature and cooling rates for each plate thickness, Figure 61. These pins each have two thermocouples in them, one which is very near the metal surface and the other that is 3mm back from the surface, Figure 62. This allows for a measurement of the heat flux, along with a measurement of metal temperature and die temperature at each thickness. Figure 63 shows the die inserts which were cut from P20 steel, for reference the size of the inserts are 7.75 inches by 7.75 inches. 146

Figure 61: Top view and side view of the Lift Specimen Die showing the location of the thermocouple pins

147

Figure 62: Drawing of the thermocouple pin inserts, one thermocouple tip is located on the surface while the other is 3mm back form the surface

Figure 63: Lift Specimen Die inserts cut from P20 steel; (Left) ejector side insert (Right) cover side insert with thermocouple pin locations

148

To produce the Lift Specimen die castings, a 250-ton die casting machine (Bühler

H-250 SC) with a Fondarex vacuum unit [Highvac Economy 1C-250] was used. This die casting machine is located in The Center for Design and Manufacturing Excellence

(CDME) at The Ohio State University (OSU). The fast shot speed was set at 1.5m/s for these experiments and was initiated as the metal reached the gating. The metal cleanliness is important to make quality HPDC casting, so the alloys were degassed with a rotary degassing unit (Pyrotek Star 2500) with argon for 20 minutes before casting. To transport the metal to the shot sleeve, a robotic ladling arm (Rimrock) was used with steel ladle that was coated with a typical boron nitride coating. Once the melt was ladled into the shot sleeve, the shot profile was initiated automatically within a second of pouring. The vacuum system was automatically turned on and off based on the position of the shot piston and calculated filling time.

5.2.3: Results and Discussion

After production, all of the plates and tensile bars were checked for porosity and defects with x-ray inspection, which was conducted at Boeing by a certified level 3 technician. The samples were rated on porosity, shrinkage, and cold fill based on ASTM

E505, which uses a rating system that goes from ‘0’ (best quality) to ‘4’ (major defects).

To study how the vacuum level affected the quality and mechanical properties of the

Lift380 alloy, samples were produced at 3 different vacuum levels and investigated. The vacuum level was changed by slowing down the slow shot speed, thus allowing for more

149 time for the system to pull vacuum. The 3 vacuum levels in this study were 85mBar,

145mBar, and 165mBar. While none of the vacuum levels were high enough to be considered super vacuum, the 85mBar is near a super vacuum level while the 165mBar is closer to a typical vacuum assisted level. Table 18 shows the percentage of tensile samples that had a particular porosity rating for each vacuum level. It can be observed that as the cavity pressure decreased (better vacuum), the higher percentage of quality tensile bar castings (‘0’ or ‘1’ ratings) are produced. The same trend can be seen with the

3mm thick bend plates in Table 19. Another observation that was made, was that the tensile bar and 5mm plate had near 100% fill rate, while the fill rate decreased with decreasing plate thickness. This was due to early solidification in the thinner plates due to the higher cooling rates. The 3mm plates were found to only successfully fill 80% of the time while the 2mm plates only filled 35% of the time. The amount of successful fills for each plate thickness was observed to improve with decreasing cavity pressure (better vacuum). This supports the conclusion that applying a vacuum helps the filling of casting by reducing the back pressure in the die cavity.

Table 18: X-Ray porosity results for tensile bars produced at different vacuum levels

Tensile Bar E505 Porosity Ratings Vacuum Level (mBar) Rating 85 145 165 0 59% 10% 5% 1 41% 62% 58% 2 0% 29% 32% 3 0% 0% 0% 4 0% 0% 0% # samples 32 21 19 150

Table 19: X-Ray porosity results for 3mm plates produced at different vacuum levels

3mm Bar E505 Porosity Ratings Vacuum Level (mBar) Rating 85 145 165 0 58% 29% 26% 1 16% 43% 37% 2 23% 14% 32% 3 0% 0% 0% 4 3% 7% 5% # samples 31 14 19

After completion of x-ray inspection on all of the tensile specimen produced at the 3 different vacuum levels, an average rating was calculated for each vacuum level. It was found that tensile bars produced with an 85mBar vacuum level had an average rating near ‘0’ (best), while bars produced at 145/165mBar had an average rating of around ‘1’.

Tensile specimen with a rating of 0 were selected for testing of mechanical properties for the 85mBar vacuum level, while samples with a rating of ‘1’ were selected to evaluate the mechanical properties of the 145/165mBar vacuum level. There were also tensile tests performed on samples produced at the 165mBar rating which had a porosity rating of ‘2’, so that a further understanding on how the porosity level affected mechanical properties could be obtained. Figure 64 shows the typical tensile curves for each vacuum level and

Table 20 presents the mechanical properties. The first thing to observe is that the x-ray porosity rating corresponds with the elongation of the samples, as the less porosity present in the specimen, the better the elongation. This is because the porosity acts as

151 initiation sites for cracking as well as helps the propagation of cracks through the sample

[100–103]. The samples produced at an 85mBar vacuum level had the least porosity and thus the best elongation. Tensile bars that were rated a ‘1’ on the porosity scale had elongations worse than ratings of ‘0’, however were found to have elongations about the same when produced at both 165mBar and 145mBar. In conclusion the introduction of vacuum to the HPDC process lead to reduction in porosity in the 3 point bend plates and tensile specimen. As the vacuum level improved, the reduction of porosity on average was greater and lead to an improvement in elongation as there were less pores to propagate cracks through the samples.

Figure 64: Typically tensile curves for Lift380 tensile specimen produced at different vacuum levels

152

Table 20: Mechanical properties for Lift380 produced at different vacuum levels

Vac E505 YS (MPa) UTS (MPa) Elongation (%) Level Rating 85 0 161.1 ± 3.0 297.3 ± 18.0 3.8 ± 0.8 145 1 171.6 ± 2.8 274.7 ± 6.7 2.5 ± 0.2 165 1 174.8 ± 5.0 278.6 ± 7.5 2.4 ± 0.1 165 2 174.4 ± 1.7 240.4 ± 0.4 1.6 ± 0.1

The incorporation of thermocouple pins in the die allowed for the cooling rates of each different plate thickness to be investigated. Figure 65 shows the cooling curves that were measured during a typical die casting cycle. It can be observed that the thicker plates have a slower cooling rate then the thinner planes, as the slope of the curve for the

5mm plate is less steep than that of the 3mm and 2mm curves for the exposed thermocouple (near surface). It can also be observed that for the internal thermocouples, the thicker plates report a higher temperature initially after casting compared to 2/3mm plates because the plate stays at a higher temperature for a longer amount of time. The exposed and internal thermocouple also takes longer to converge for the thicker plates.

All the thermocouples eventually converge within a few degrees before the next cycle begins which happens 70-90 seconds later. The actual cooling rates of each thickness were not able to be calculated because the maximum temperatures measured were short of liquidus temperatures. This is because the thermocouple response time is not quick enough for the high velocity of the molten metal and fast cooling rates. Further development in the thermocouples to increase the response time is needed to measure at these cooling rates.

153

Figure 65: Measured cooling curves for each plate thickness

The grain sizes of the different thickness plates were measured by using electron backscatter diffraction (EBSD) techniques. Figure 66 shows the grain size distribution for the 2mm plate, which had an average grain size of 18μm in the center of the plate where the cooling rate is the slowest. Near the edge of the plate, where the cooling rate is the quickest, the grain size was measured to be much smaller with an average size of 5.3μm.

When compared to the 5mm plate, Figure 67, the difference in grain size was obvious in the center of the plate, where the average grain size was 25μm. However, at the edge of the plate the different in grain size was not as large as the 5mm plate had an average grain size of 5.5μm. This grain size difference was reduced at the edge because of how high the cooling rates are for both samples near the edge. Because of the large difference in grain

154 size in the bulk (center) of the plants, the plates will have different yield strengths as described by the Hall-Petch relationship. The skin (or edge) of the casting will have an increased yield strength compared to the bulk yield strength as well. It is important to determine the contribution each of these regions have to the overall yield strength of the plate. The contribution should scale with the thickness as the skin yield strength contribution should be more apparent in the 2mm sample compared to the 5mm sample as the skin is a larger percentage of the casting. Studies are being performed to determine the yield strength both the skin and bulk areas. If the strength contribution from each of these regions is able to be determined and related to thickness, more accurate yield strength prediction can be calculated for areas of different thickness. These strengths can then be related to the grain size measured for each region, which can then be related to the solidification rate that can be calculated for each region. Finally a relationship between the simulated cooling rates and the calculated solidification rates can be established. This will enable the yield strengths of different thickness areas to be predicted based solely on the simulated cooling rates and thickness. This will allow for casting designers to unlock further weight savings by designing components which minimized thicknesses in specific areas which still meet the minimum mechanical properties required. With the castings produced and data obtained from the Lift specimen die, it will now be possible determine these relationships and mechanical properties; further research is ongoing to complete this investigation.

155

Figure 66: Electron backscatter diffraction analysis (EBSD) and grain size distribution of the 2mm plate; alloy: EZCast

Figure 67: Electron backscatter diffraction analysis (EBSD) and grain size distribution of the 5mm plate; alloy: EZCast

156

5.2.4: CANMET industry application

To further investigate the super vacuum die casting process, an industrial component (side impact bar) was cast at CANMET (Hamilton, Canada) and investigated.

The side impact bar (SIB) measures about 30 inches in length by 3 inches in width and had an average thickness of 3mm throughout the majority of the casting, Figure 68. These castings were produced with both the EZCast and Lift380 alloys. The melt temperature of each alloy was set to about 115°C above the liquidus, while the die temperature was kept at 120°C. The fast shot speed was 4.5m/s for the casting of Lift380, while 4m/s was used for EZCast. The EZCast was found to fill slightly better, because of its lower viscosity, so a lower fast shot speed could be used to make quality castings. These components were cast both with the super vacuum die casting system operational and with it turned off

(conventional HPDC). The die cavity has an average pressure of 8mBar just before the metal entered the cavity in the shots that the super vacuum system was activated.

157

Figure 68: Drawing of side impact bar (SIB) with overall dimensions

Simulation using MAGMASOFT was performed to find initial processing parameters and to predict the formation of porosity in the samples. After the casting campaign, simulations were ran again with the casting parameters that were optimized during the casting trial and used to produce the components. Simulations showed that no formation of porosity was predicted for the EZCast castings, while a small percentage of porosity was predicted to form in the castings produced with the LIFT380 alloy, Figure

69. This porosity was concentrated around the ribbing in the SIB, which were predicted hot spots. Lift380 has a larger solidification range compared to EZCast and thus is more susceptible to the formation of shrinkage porosity. Lift380 has a solidification range of

109°C, while EZCast has a range half as long at only 52°C, Table 15.

158

Figure 69: MAMGASOFT simulation predicting porosity formation in SIB castings for both (Top) EZCast and (Bottom) Lift380 alloys

159

After the components were produced at CANMET, they were sent to Boeing for x-ray inspection, Figure 70. The x-ray inspection was performed by a certified level 3 technician and ATSM E505 was used for rating the castings. Both the conventional

HPDC components and SVDC components were inspected for each alloy. Figure 71, shows x-ray results for the EZCast alloy. The part on the bottom was produced with poor vacuum (730mbar) and showed a large amount of porosity. This was the trend found for many samples that were produced without vacuum or with poor vacuum conditions. On the top is a typical sample produced with a vacuum between 7-10mBar. There were no indications in the majority of the castings, while some casting showed limited indications of porosity. This confirms that the SVDC process had a positive effect on decreasing the amount of entrapped air porosity by reducing the air in the cavity before filling.

Figure 70: Picture of the SIB casting produced at CAMMET with gating and overflows still attached

160

The same comparisons were performed for the Lift380 alloy, Figure 72. When observing Lift380 SIB parts that were produced without vacuum (1010mBar), a large amount of porosity was observed in the x-ray results. This matches the results that were found in the conventional HPDC EZCast samples, once again supporting the SVDC reduced the amount of entrapped air porosity. For the majority of Lift380 SIB parts produced with vacuum (7-10mBar), a small amount of porosity was detected. It is important that even though these components were never free of porosity, they did still meet the requirements to be considered a quality casting to be used in commercial applications. This porosity which was observed is predicted to be solidification porosity, as it was found mostly near the ribs which were predicted as hot spot areas and areas prone to form porosity in the simulations. This was only observed in the Lift380 alloy simulations, as it has a larger solidification range, and was only observed in the Lift380 castings via x-ray inspect. The EZCast parts showed very little solidification shrinkage in actual castings while this matched the simulation prediction.

161

Figure 71: X-ray results of SIB castings produced with the EZCast alloy

Figure 72: X-ray results of SIB castings produced with the Lift380 alloy

162

5.2.5: Conclusions

The design of the Lift Specimen die allowed for the cooling rate and vacuum level effect to be studied on varying plate thicknesses. When using a vacuum assisted casting process, a higher percentage of thin-wall castings were able to achieve complete filling.

All of the plates and tensile bars produced were investigated with x-ray inspection techniques to determine the quantity of porosity in each produced at different vacuum levels. It was observed that castings which were produced at higher vacuum levels (better vacuum, lower pressure) had less porosity on average. The tensile bars were produced at three different vacuum levels, 85mBar, 145mBar, and 165mBar to study how the vacuum affects the mechanical properties. The better vacuum levels produced samples that had over twice the elongation which also resulted in higher ultimate tensile strengths.

However the yield strength did not change significantly with tensile bars produced at different vacuum levels.

Simulated and measured cooling rates were found for each of the plate thicknesses and it was observed that decreasing thickness resulted in higher cooling rates and thus higher solidification rates. Because of these differences in cooling rates, the grain size in the plates varied greatly with thickness. The 5mm plate had an average grain size of 25μm, while the 2mm plate only had an average grain size of 18μm. At the edge of the plates, referred to as the skin region, the grain size was further reduced for each thickness because of the increase in cooling rate in this region. Each of these regions will have different mechanical properties because of the grain size difference as described by

163 the Hall-Petch equation. It is important to determine the contribution that each of these regions have on the overall yield strength of the plates. If the strength contribution from each of these regions is able to be determined and related to thickness, more accurate yield strength predictions can be calculated for areas of different thickness. These strengths can then be related to the grain size measured for each region, which can be related to the solidification rate that can be calculated for each region. A final relationship between the simulated cooling rates and the calculated solidification rates can be established so that the yield strengths of different thickness areas can be calculated based solely on the simulated cooling rates. With this information, designers will have a better idea of strengths for different thickness sections in castings, allowing casting designs to be optimized based on the location specific mechanical properties needed.

The SIB casting trials showed good agreement with the Lift Specimen die vacuum studies, as it demonstrated that the use of the super vacuum process lead to a large reduction in porosity from gas entrapment. Solidification porosity was still observed in

SIB parts produced with the Lift380 alloy, as the solidification range for this alloy is very large. However in the EZCast alloy which had a smaller solidification range this shrinkage porosity was not as commonly observed. These findings were compared with casting simulation software predictions, which also showed porosity formation in Lift380 but not EZCast. Further evaluation of the mechanical properties of these castings produced with and without vacuum is needed to be performed. Location specific studies also need to be conducted on these castings to study the grain size, as well as the mechanical properties in different areas of the castings.

164

Chapter 6: Summary and Future Work

6.1: Summary

This dissertation summarized the design and development of two new magnesium alloys and the process development of the super vacuum die casting process. These developments are important so that further light weighting of automobiles and other technologies can be achieved to lessen environmental impact, lower costs, and to reduce energy consumption. The CALPHAD (CALculation of PHAse Diagrams) method was used in development of AT and ATS alloys and to aid in the design of heat treatment schedules. It was showed to be an important aid to expedite the development of new tailored alloy systems.

The results on Mg-Al-Sn alloy development have shown the addition of Sn has a positive influence on mechanical properties. The as-cast microstructure of this alloy consists of Mg, Mg2Sn, and Mg17Al12 which were concentrated around the gain boundaries and provided strengthening. Artificially aging treatments (T5) lead to the precipitation of fine Mg17Al12 and Mg2Sn precipitates however this treatment was not very effective in increasing the mechanical properties and had a slow response time.

Solution treatments designed with use of the CALPHAD method lead to the dissolution

165 of Mg17Al12, however only partial dissolution of Mg2Sn was achieved because of its higher thermal stability and lower solubility. Mechanical testing revealed that as-cast

AT72 possesses an increase of YS compared to AM60B while only having a slightly lower elongation though these strengths were still short of AZ91. Artificial aging of

AT72 led to a further increase in YS, however a large reduction of elongation was found due to this heat treatment which produced undesirable mechanical properties.

The investigation of Mg-Al-Sn-Si alloys followed AT72 development and demonstrated that the addition of Si had a further positive influence on mechanical properties. The as-cast microstructure of the ATS alloy produced by gravity casting consisted of Mg, Mg17Al12, Mg2Sn, and Mg2Si. As was observed in AT72, the secondary phases aggregated near the gain boundaries and provide strengthening to the alloy. The introduction of Si leads to an additional Mg2Si phase forming which provided further increased strengths. Once again the T5 artificially aging treatments were not very effective on the ATS alloy, however were slightly improved when compared to AT72 because of the introduction of Si which aided in the precipitation of Mg2Sn and Mg17Al12.

In attempt to improve the dissolution of intermetallics during the solution treatment, a novel two stage treatment was designed with use of the CALPHAD method. The two stage treatment, which consisted of a long lower temperature hold paired with a shorter high temperature hold, leads to further dissolution of the Mg2Sn phase.

The as-cast mechanical properties of ATS provided an increase to YS and UTS when compared with AT72 credited to the introduction of the Mg2Si intermetallics.

Various heat treatments were explored, however increased mechanical properties were

166 not found in any treatment that involved solution treatment because of the blistering of porosity that occurred at higher temperatures. The T5 condition showed a slight increase in YS and provided the best combination of properties for the ATS alloy when compared to current magnesium alloys.

To increase mass savings in lightweight applications, thin-wall castings are becoming larger and further process optimization and validation is important to insure quality castings. Two dies were developed to study both the fluidity of alloys and the effects that vacuum die casting has on the filling of thin wall sections and mechanical properties. The fluidity die shows how different HPDC processing parameters can affect the flow length of the alloys. Input parameters such as metal specific properties, die properties, and processing parameters are able to be generated for specific alloys and used to improved casting simulations. In these experiments, flow simulations using alloy- specific properties (viscosity, thermophysical data, etc.) for both EZCast and Lift380 alloys were shown to predict the flow length and flow pattern with reasonable accuracy, especially the trend of flow length that 3mm > 2mm > 1mm passage at the same time step. With the development of this fluidity die, HPDC processing parameters can be studied and accurately compared to simulations.

The Lift Specimen die aided in the investigation of super vacuum die casting optimization for thin-wall castings and validated that vacuum die casting leads to improvement mechanical properties. With the absence of air in the die cavity, the filling of castings improved and for formation of porosity was reduced, leading to increased mechanical properties. The SIB casting trials showed good agreement with Lift Specimen

167 die vacuum studies, as it demonstrated that the introduction of super vacuum to the die casting process lead to a large reduction in porosity which previously formed due to gas entrapment. Shrinkage porosity was observed in the Lift380 castings as the solidification range for this alloy is very large and makes shrinkage porosity formation more likely.

However in the EZCast alloy, which had a smaller solidification range, this shrinkage porosity was not as commonly observed. The casting simulation software predicted these same results, as it predicted porosity formation in Lift380 but not EZCast.

6.2: Future work

The precipitation of Mg2Sn and Mg17Al12 needs to be investigated more thoroughly so a complete understanding of the orientation and fraction percent of these precipitates can be obtained. The orientation of these precipitates is very important, as certain orientations contribute to strengthening more than others. Measurements of the fraction percent of the each of the precipices could provide further insight on how the Si addition and heat treatment schedule effects the precipitation of each phase. Furthermore micro alloying additions, such as Na and Zn, should studied to determine if they can be used to shorten treatments times or improve the response as they have shown promising results in previous Mg-Sn alloys [63,68].

For the process development, further investigation of the different thickness plates needs to be performed so that the grain size of the bulk and skin can be related to the mechanical properties of each plate. With the measured contributions of strength from the

168 core and skin locations, specific mechanical properties can be predicted for varying thicknesses throughout a casting. With this information, simulated cooling rates can be used to give designers a better idea of strengths for different thickness sections in castings and will allow for designs to be optimized for location specific mechanical properties.

Further evaluation of the mechanical properties of the SIB castings produced should be used to validate the location specific properties found from calculation and simulation.

169

References

[1] Ribeiro K, Kobayashi S, Beuthe M, Gasca J, Greene D, Lee DS, et al. Transport

and its infrastructure. In: Metz B, Davidson O., Bosch PR, Dave R, Meyer L.,

editors. Clim. Chang. 2007 Mitig., 5: Cambridge University Press; 2007.

[2] Joost W, Davis P, Eberhardt J, Bezanson C, Patterson P, Wicker J, et al. Energy,

Materials, and Vehicle Weight Reduction. 2012.

[3] Casadei A, Broda R. Impact of Vehicle Weight Reduction on Fuel Economy for

Various Vehicle Architectures. 2008.

[4] King JF. Magnesium: commodity or exotic? Mater Sci Technol 2007;23:1–14.

doi:10.1179/174328407X154374.

[5] Pekguleryuz M, Celikin M. Creep resistance in magnesium alloys. Int Mater Rev

2010;55:197–217. doi:10.1179/095066010X12646898728327.

[6] Pekguleryuz MO, Kaya AA. Creep Resistant Magnesium Alloys for Powertrain

Applications. Adv Eng Mater 2003;5:866–78. doi:10.1002/adem.200300403.

[7] Hort N, Huang Y, Kainer KU. Intermetallics in Magnesium Alloys. Adv Eng

Mater 2006;8:235–40. doi:10.1002/adem.200500202.

[8] Avedesian, M.M.; Baker H. Magnesium and Magnesium Alloys. Materials Park,

OH: ASM International; 1999.

[9] Cheng J, Ghosh S. A crystal plasticity FE model for deformation with twin

170

nucleation in magnesium alloys. Int J Plast 2015;67:148–70.

doi:10.1016/j.ijplas.2014.10.005.

[10] Penghuai F, Liming P, Haiyan J, Wenjiang D, Chunquan Z. Tensile properties of

high strength cast Mg alloys at room temperature : A review. China Foundry

2014;11:277–86.

[11] Mordike BL, Lukac P. Physical Metallurgy. Magnes. Technol. Metall. Des. Data,

Appl., vol. M, Springer Science & Business Media; 2006, p. 63–103.

[12] Yang K V., Cáceres CH, Easton M a. Strengthening Micromechanisms in Cold-

Chamber High-Pressure Die-Cast Mg-Al Alloys. Metall Mater Trans A

2014;45:4117–28. doi:10.1007/s11661-014-2326-x.

[13] Nagasekhar A V., Easton M a., Caiceres CH. Solute Content and the Grain

Microstructure of High Pressure Diecast Magnesium-Aluminium Alloys. Adv Eng

Mater 2009;11:912–9. doi:10.1002/adem.200900175.

[14] Labusch R. A Statistical Theory of Solid Solution Hardening. Phys Status Solidi

1970;41:659–69. doi:10.1002/pssb.19700410221.

[15] Orowan E. Discussion on Internal Stresses. Symp. Intern. Stress. Met., London:

The Insitute of Metals; 1948, p. 451–3.

[16] Nayeb-Hashemi AA, Clark JB. Phase Diagrams of Binary Magnesium Alloys.

Metals Park, Ohio: ASM International; 1988.

[17] Muddle BC, Ringer SP, Polmear IJ. High Strength Microalloyed Aluminum

Alloys. Adv. Mater. ’93, VI - A B, ELSEVIER SCIENCE PUBL B V; 1994, p.

999–1023.

171

[18] Polmear IJ. Control of precipitatikon processes and properties in aged aluminum

alloys by microalloying. Mater Forum 1999;23:117–35.

[19] Polmear IJ. Magnesium alloys and applications. Mater Sci Technol 1994;10.

[20] Luo A, Pekguleryuz MO. Review Cast magnesium alloys for elevated temperature

applications. J Mater Sci 1994;29:5259–71.

[21] Ghali E. Magnesium and Magnesium Alloys. In: Revie RW, editor. Uhlig’s

Corros. Handb. 3rd ed., John Wiley & Sons, Inc; 2011, p. 809–35.

[22] Mordike BL, Ebert T. Magnesium Properties — applications — potential. Mater

Sci Eng A 2001;302:37–45.

[23] Kulekci MK. Magnesium and its alloys applications in automotive industry. Int J

Adv Manuf Technol 2007;39:851–65. doi:10.1007/s00170-007-1279-2.

[24] Luo A a. Magnesium: Current and Potential Automotive Applications. Jom

2002;54:42–8. doi:10.1007/BF02701073.

[25] Luo AA. Wrought Magnesium Alloys and Manufacturing Processes for

Auntomotive Applications. SAE, 2005.

[26] Song J, Xiong S, Li M, Allison J. The correlation between microstructure and

mechanical properties of high-pressure die-cast AM50 alloy. J Alloys Compd

2009;477:863–9. doi:10.1016/j.jallcom.2008.11.040.

[27] Luo AA. Recent magnesium alloy development for elevated temperature

applications. Int Mater Rev 2004;49:13–30. doi:10.1179/095066004225010497.

[28] Luo a. a. Recent magnesium alloy development for elevated temperature

applications. Int Mater Rev 2004;49:13–30. doi:10.1179/095066004225010497.

172

[29] Chang YA, Chen S, Zhang F, Yan X, Xie F, Schmid-fetzer R, et al. Phase

diagram calculation : past , present and future. Prog Mater Sci 2004;49:313–45.

doi:10.1016/S0079-6425(03)00025-2.

[30] Cao W, Chen S, Zhang F, Wu K, Yang Y, Chang YA, et al. PANDAT software

with PanEngine , PanOptimizer and PanPrecipitation for multi-component phase

diagram calculation and materials property simulation. Calphad 2009;33:328–42.

doi:10.1016/j.calphad.2008.08.004.

[31] Andersson, J-O; Helander, T; Hoglund, L; Shi, P; Sundman B. THERMO-CALC

& DICTRA, Computational Tools For Materials Science J-O Andersson, Thomas

Helander,Lars Hdghmd, Pingfang Shi, Bo Sundman. Calphad 2002;26:273–312.

[32] Bale CW, Bélisle E, Chartrand P, Decterov SA, Eriksson G, Gheribi AE, et al.

FactSage thermochemical software and databases , 2010 – 2016. Calphad

2016;54:35–53. doi:10.1016/j.calphad.2016.05.002.

[33] Kaufman L, John A. CALPHAD , first and second generation – Birth of the

materials genome. Mater Sci Eng 2014;70:3–6.

doi:10.1016/j.scriptamat.2012.12.003.

[34] Chang, YA; Yang Y. Application of Computational Thermodynamics to Rapidly

Determine Multicomponent Phase Diagrams. In: Zhao J, editor. Methods phase

Diagr. Determ., Oxford, UK: Elsevier, Ltd; 2007, p. 273.

[35] Luo A a., Powell BR, Sachdev AK. Computational phase equilibria and

experimental investigation of magnesium–aluminum–calcium alloys.

Intermetallics 2012;24:22–9. doi:10.1016/j.intermet.2012.01.001.

173

[36] Twarog, D., D. Apelian AL. High Integrity Casting of Lightweight Components.

Arlington Heights, Illinois: North American Die Casting Association; 2016.

[37] Luo A a. Magnesium casting technology for structural applications. J Magnes

Alloy 2013;1:2–22. doi:10.1016/j.jma.2013.02.002.

[38] Vinarcik EJ. High Integrity Die Casting Processes. New York City, NY: John

Wiley & Sons, Inc; 2003.

[39] Cashion SP. The mechanism of protection of molten magnesium by cover gas

mixtures containing sulphur hexafluoride. J Light Met 2 2002;2:43–7.

[40] Lyon P, Ricketts NJ. Magnesium melt protection at Magnesium Elektron using

HFC-134a. Magnes. Technol., 2003.

[41] Pekguleryuz MO, Dieringa H, Bohlen J, Hort N, Letzig D, Kainer KU. Advances

in Manufacturing Processes for Magnesium Alloys. Magnes Technol 2010:19–24.

[42] Ricketts NJ, Cashion SP. Hydrofluorocarbons as a Replacement for Sulphur

Hexafluoride in Magnesium Processing n.d.:1–7.

[43] Emley, EF; Fisher P. The Control of Quality of Magnesium-Based Alloy

Castings. J Institue Met 1957;85.

[44] Baker W. Microporosity in magnesium alloy castings. J Institue Met 1945;71:165.

[45] Brown, Z; Luo, AA;Sadayappan, K; Ouimen, LJ;Zindel, J; Beals RM.

Development of super-vacuum die casting process for mangesium alloys, North

American Die Casting Association Transactions; 2009, p. T09-43.

[46] Niu XP, Hu BH, Pinwill I, Li H. Vacuum assisted high pressure die casting of

aluminium alloys. J Mater Process Technol 2000;105:119–27. doi:10.1016/S0924-

174

0136(00)00545-8.

[47] Chill Vents. Castool Tool Syst n.d. www.castool.com (accessed January 11, 2017).

[48] Bowles AL, Blawert C, Hort N, Kainer KU. Microstructural Investigation of the

Mg-Sn and Mg-Sn-Al Alloy Systems. Magnes Technol 2004:307–10.

[49] Jung IH, Kang DH, Park WJ, Kim NJ. Mg-Sn Phase Diagram. Int J Mater Res

2007.

[50] Hort N, Huang Y, Abu Leil T, Maier P, Kainer KU. Microstructural Investigations

of the Mg-Sn-xCa System. Adv Eng Mater 2006;8:359–64.

doi:10.1002/adem.200600014.

[51] Liu H, Chen Y, Tang Y, Wei S, Niu G. The microstructure, tensile properties, and

creep behavior of as-cast Mg–(1–10)%Sn alloys. J Alloys Compd 2007;440:122–

6. doi:10.1016/j.jallcom.2006.09.024.

[52] Luo A a., Fu P, Peng L, Kang X, Li Z, Zhu T. Solidification Microstructure and

Mechanical Properties of Cast Magnesium-Aluminum-Tin Alloys. Metall Mater

Trans A 2011;43:360–8. doi:10.1007/s11661-011-0820-y.

[53] Nie J-F. Precipitation and Hardening in Magnesium Alloys. Metall Mater Trans A

2012;43:3891–939. doi:10.1007/s11661-012-1217-2.

[54] Celotto S. TEM study of continuous precipitation in Mg–9 wt%Al–1 wt%Zn

alloy. Acta Mater 2000;48:1775–87. doi:10.1016/S1359-6454(00)00004-5.

[55] ASTM Standard B557-15. Standard Test Methods for Tension Testing Wrought

and Cast Aluminum and Magnesium Alloy Products. ASTM Int., West

Conshohocken, PA: 2016. doi:10.1520/B0557-15.2.

175

[56] Pandat software package and PanAl thermodyamic database 2017.

[57] Fisher K. Fundamentals of Solidification. 4th Editio. CRC Press; 1998.

[58] Sun W, Shi X, Cinkilic E, Luo AA. Investigation of the non-equilibrium

solidification microstructure of a Mg – 4Al – 2RE ( AE42 ) alloy. J Mater Sci

2016;51:6287–94. doi:10.1007/s10853-016-9925-4.

[59] Berman TD, Li M, Allison JE. Microsegregation in High Pressure Die Cast Mg

Alloys. Magnes Technol 2017:73–8. doi:10.1007/978-3-319-52392-7.

[60] Luo AA, Fu P, Peng L, Kang X, Li Z. Solidification Microstructure and

Mechanical Properties of Cast Magnesium-Aluminum-Tin Alloys 2013;43.

doi:10.1007/s11661-011-0820-y.

[61] Larsson H. A Model for 1D Multiphase Moving Phase Boundary Simulations

under Local Equilibrium Conditions. Calphad 2014;47:1–8.

[62] Nie JF. Effects of precipitate shape and orientation on dispersion strengthening in

magnesium alloys. Scr Mater 2003;48:1009–15. doi:10.1016/S1359-

6462(02)00497-9.

[63] Sasaki TT, Oh-ishi K, Ohkubo T, Hono K. Enhanced age hardening response by

the addition of Zn in Mg–Sn alloys. Scr Mater 2006;55:251–4.

doi:10.1016/j.scriptamat.2006.04.005.

[64] Gibson M a., Fang X, Bettles CJ, Hutchinson CR. The effect of precipitate state

on the creep resistance of Mg–Sn alloys. Scr Mater 2010;63:899–902.

doi:10.1016/j.scriptamat.2010.07.002.

[65] Mendis CL, Bettles CJ, Gibson M a., Gorsse S, Hutchinson CR. Refinement of

176

precipitate distributions in an age-hardenable Mg–Sn alloy through microalloying.

Philos Mag Lett 2006;86:443–56. doi:10.1080/09500830600871186.

[66] Shi X, Li D, Luo A a., Hu B, Li L, Zeng X, et al. Microstructure and Mechanical

Properties of Mg-7Al-2Sn Alloy Processed by Super Vacuum Die-Casting. Metall

Mater Trans A 2013;44:4788–99. doi:10.1007/s11661-013-1799-3.

[67] Braszczynaska-Malik KN. Discontinuous and continuous precipitation in

magnesium – aluminium type alloys. J Alloys Compd 2009;477:870–6.

doi:10.1016/j.jallcom.2008.11.008.

[68] Wang X, Gibson M, Wang XJ, Zhu SM, Easton MA, Gibson MA, et al. Heat

treatment of vacuum high pressure die cast magnesium alloy AZ91 Heat treatment

of vacuum high pressure die cast magnesium alloy AZ91. Int J Cast Met Res

2014;27:161–6. doi:10.1179/1743133613Y.0000000091.

[69] Van Der Planken J. Precipitation Hardening in Magnesium-Tin Alloys. J Mater

Sci 1969;4:927–9. doi:10.1007/BF00549785.

[70] Mendis CL, Bettles CJ, Gibson M a., Hutchinson CR. An enhanced age hardening

response in Mg–Sn based alloys containing Zn. Mater Sci Eng A 2006;435–

436:163–71. doi:10.1016/j.msea.2006.07.090.

[71] Lee SG, Jeon JJ, Park KC, Park YH, Park IM. Investigation on microstructure and

creep properties of as-cast and aging-treated Mg–6Sn–5Al–2Si alloy. Mater Sci

Eng A 2011;528:5394–9. doi:10.1016/j.msea.2011.03.080.

[72] Kang DH, Park SS, Kim NJ. Development of creep resistant die cast Mg–Sn–Al–

Si alloy. Mater Sci Eng A 2005;413–414:555–60. doi:10.1016/j.msea.2005.09.022.

177

[73] Li X, Chen Y, Wang X, Ma G. Effect of cooling rates on as-cast microstructures

of Mg-9Al-xSi (x=1, 3) alloys. Trans Nonferrous Met Soc China 2010;20:s393–6.

doi:10.1016/S1003-6326(10)60504-3.

[74] Beer, S; Frommeyer G. Development of Mg2Si Light Weight Alloys. Magnes

Alloy Their Appl 1992:317–24.

[75] Du J, Iwai K, Li W, Peng J. Effects of alternating current imposition and alkaline

earth elements on modification of primary Mg2Si crystals in hypereutectic Mg-Si

alloy. Trans Nonferrous Met Soc China 2009;19:1051–6. doi:10.1016/S1003-

6326(08)60405-7.

[76] Wang H, Wang W, Zha M, Zheng N, Gu Z, Li D, et al. Influence of the amount of

KBF4 on the morphology of Mg2Si in Mg – 5Si alloys. Mater Chem Phys

2008;108:353–8. doi:10.1016/j.matchemphys.2007.10.006.

[77] Pandat software package and PanMagnesium thermodyamic database 2017.

[78] Cáceres CH, Poole WJ, Bowles a. L, Davidson CJ. Section thickness,

macrohardness and yield strength in high-pressure diecast magnesium alloy AZ91.

Mater Sci Eng A 2005;402:269–77. doi:10.1016/j.msea.2005.04.024.

[79] Aghion E, Moscovitch N, Arnon A. The correlation between wall thickness and

properties of HPDC Magnesium alloys. Mater Sci Eng A 2007;447:341–6.

doi:10.1016/j.msea.2006.10.065.

[80] Campbell J. Incipient grain boundary melting. Mater Sci Technol 2008;0:1–2.

doi:10.1179/174328408X311080.

[81] Lu Y, Wang Q, Zeng X, Ding W, Zhu Y. Effect of Si on the precipitation

178

behavior of Mg-6Al alloy. J Mater Sci Lett 2001;20:397–9.

[82] Duly, D; Brechet, Y; Chenal B. Macroscopic Kinetics of Discontinuous

Precipitation in a Mg-8.5 wt% Al Alloy. Acta Mater 1992;40:2289–300.

[83] Hornbogen E. Systematics of the Cellular Precipitation Reactions. Metall Trans

1972;3.

[84] Timelli, G; Otto, L; Arnberg, L; Laukli H. Effect of Solution Heat Treatments on

the Microstructure and Mechanical Properties of a Die-Cast AlSi7MgMn Alloy.

Metall Mater Trans 2008;39A:1747–58. doi:10.1007/s11661-008-9527-0.

[85] Nehan M, Maloney R. Magnesium AM60B Instrument Panel Structure for

Crashworthiness FMVSS 204 and 208. SAE Pap. No. 960419, Warrendale, PA:

1996.

[86] Gibbs S. Magnesium structural part parts with Myth. Met Cast Des Purch

2010:29–33.

[87] Riopelle L. Magnesium Applications. Int. Magnes. Assoc. Annu. Magnes.

Automot. Semin., Livonia, MI: 2004.

[88] Sabatino M Di, Arnberg L, Brusethaug S, Apelian D. Fluidity evaluation methods

for Al – Mg – Si alloys. Int J Cast Met Res 2006;19:94–7.

doi:10.1179/136404606225023345.

[89] Sabatino M Di, Sahnkar S, Apelian D, Arnberg L. Influence of temperature and

alloying elements on fluidity of Al-Si Alloys. In: Tiryakioglu M, Crepeau PN,

editors. Shape Cast. John Campbell Symp., TMS (The Minerals, Metals &

Materials Society); 2005.

179

[90] Çolak M, Kayikci R, Dispinar D. Influence of Different Cross Sections on

Fluidity Characteristics of A356. Trans Indian Inst Met 2015;68:275–81.

doi:10.1007/s12666-014-0455-y.

[91] Zhao P, Wang Q, Liu L, Wei Z, Zhai C. Fluidity of Mg-Al-Ca alloys in the high-

presure die casting process. Int J Mater Res 2007;98.

[92] Kim Y-C, Choi S-W, Kim C-W, Cho J-I, Kang C-S. Influence of process

parameters on the fluidity of high pressure die-casting Al-Si Alloys. Adv Mater

Res 2013;813:171–4.

[93] Han Q, Xu H. Fluidity of alloys under high pressure die casting conditions. Scr

Mater 2005;53:7–10.

[94] Ravi B. Casting Simulation – Best Practices. Trans. 58th IFC, Ahmedabad: 2009.

[95] ProCAST 2017.

[96] MAGMASOFT 2017.

[97] Flender E, Sturm J. Technical Development Report - Thirty years of Casting

Process Simulation. Int J Met 2010:7–23.

[98] ProCAST 2016.1 User’s Guide. ESI Group; 2017.

[99] Guo Z-P, Xiong S-M, Liu BC, Li M, Allison J. Development of thermal boundary

condition model at metal/die interface of high pressure die casting process. Int J

Cast Met Res 2011;24:151–7. doi:10.1179/136404611X13001912813825.

[100] Li X, Xiong SM, Guo Z. Journal of Materials Science & Technology Correlation

between Porosity and Fracture Mechanism in High Pressure Die Casting of

AM60B Alloy. J Mater Sci Technol 2016;32:54–61.

180

doi:10.1016/j.jmst.2015.10.002.

[101] Chadha G, Allison JE, Jones JW. The Role of Microstructure on Ductility of Die-

Cast AM50 and AM60 Magnesium Alloys. Metall Mater Trans A 2007;38A:286–

97. doi:10.1007/s11661-006-9027-z.

[102] Lee SG, Patel GR, Gokhale AM, Sreeranganathan A, Horstemeyer MF.

Quantitative fractographic analysis of variability in the tensile ductility of high-

pressure die-cast AE44 Mg-alloy. Mater Sci Eng A 2006;427:255–62.

doi:10.1016/j.msea.2006.04.108.

[103] Weiler JP, Wood JT, Klassen RJ, Maire E, Berkmortel R, Wang G. Relationship

between internal porosity and fracture strength of die-cast magnesium AM60B

alloy. Mater Sci Eng A 2005;395:315–22. doi:10.1016/j.msea.2004.12.042.

[104] Uchida M. Development of vacuum die-casting process. China Foundry

n.d.;6:137–44.

181