A NOVEL APPROACH FOR PROPERTY MODIFICATION OF CAST ALUMINUM ALLOYS WITH NANOSTRUCTURED CHEM IC AL ADDITIONS By Yang Lu A DISSERTATION Submitted to Michigan State University in partial fulfillment of the requirements for the degree of Materials Science and Engineering - Doctor of Philosophy 201 9 ABSTRACT A NOVEL APPROACH FOR PROPERTY MODIFICATION OF CAST ALUMINUM ALLOYS WITH NANOSTRUCTURED CHEMI C AL ADDITIONS By Yang Lu Continued weight reduction which involves replacing st eel components with lightweight materials is one of the major strategies to achieve the average fuel economy standards in the automotive industry . Based on the significant success of aluminum alloys in non - structural components, aluminum alloys in structural applications attract particular interest. To this end, the ultimate objective is to achieve aluminum alloys w ith high ductility and strength. However, conventional aluminum - silico n (Al - Si) based casting alloys possess low ductility as compensation for the required high strength, since the presence of irregular eutectic silicon crystals in the as - cast condition promotes crack initiation and propagation. Strontium (Sr) is widely used to modify the Si eutectic in commercially available "high ductility" cast aluminum alloys. However, fading issues arise due to Sr oxidation at the casting temperature, and become major concern s for aluminum foundries . Therefore, it is important to develop a univer sal modifier that can be used in alloys with a larger range of Si content. In this thesis, t risilanol p olyhedral silsesquioxane (TSP) was first incorporated in Sn - based binary and ternary alloys , leading to microstructure refinement and high cree p resistance. Based on these preliminary results, TSP was further applied to Al based alloys. First, TSP was coated on Al - 12Si powders using dip - coating approach . After melting, t he microstructure of both primary Al and eutectic microconstituent s of Al - 12Si ingots was r efined . Ductility increased from 5% with no TSP treatment to 18% when treated with TSP. More interestingly, the modified structure was maintained with 150 ppm of TSP addition . The incorporation of TSP reduced the undercooling and arrest temperatures of Al - Si eutectic during solidification. These results suggest TSP bonds with Al to slow down the Al segregation from Al - Si melt during eutectic reaction, leading to the microstructural refinement of Al - Si eutectic microconstituents . Based o n the laboratory - scale results , a new TSP master was produced to facilitate the incorporation of TSP in to traditional foundry process . The optimized master composition contained 6%TSP in an Al - 12Si alloy , and 10% Al - 12Si - 6TSP master was applied to modify A l - Si binary and commercial alloys at a 100 lbs. scale. In Al - Si binary alloys, TSP addition results 50% improvement in ductility over the Al - 7.5Si base alloy without sacrificing the strength. The microstructure of both primary Al and eutectic microconstit uent s of the Al - 7.5Si alloy with TSP addition was refined. Minimal fade was observed up to a 192 - hour hold at 720 ° C , which shed s light on the Sr fading issue . In Aural TM 2 and W319 - type commercial alloys, TSP addition s reduced the solidification shrinkage of castings. M etallography showed refined Si in the Al - Si eutectic micro constituent and reduced Al secondary dendrite arm spacing (SDAS) , resulting in an 10 0% improvement in ductility without sacrificing strength . Besi des microstructural refinement, (Al 2 Cu) increase d while Q ( Al 4 Cu 2 Mg 8 Si 7 ) decreased with TSP addition, suggesting TSP changes the Cu intermetallic compounds (IMCs) demonstrated its potent ial as a strengthen ing precipitate. This study address es fundamental questi ons on how the addition of TSP influence s microstructure and s olidification behavior of Al - Si based casting alloys , thus provid ing solutions to reinforce aluminum alloys in body structure s for the environmental and fuel economy benefits . iv This thesis is dedicated to the memory of my father, Xuhui Lu, who would have been happy to see me finish my Ph.D. degree. It is also dedicated to my mother, Xiuyun Mi, my wife, Chun Liu, and my daughter, Danica Lu, for their love and support. v ACKNOWLEDGEMENTS The first person I would like to thank is my advisor, Prof. Andre Lee. I would like to expre ss my sincere gratitude for his inspiring guidance and constant encouragement during my PhD studies. His knowledge, creativeness, and passion for research have not only earned him a reputation but also position ed him as a role model for his students, inclu ding me. I would also like to thank my committee members Prof. Martin Crimp, Prof. Yue Qi, and Prof. Patrick Kwon for their intelligent questions and feedback for my research. I have been extremely lucky to have Larry Godlewski, Jacob Zindel, and Mei Li who cared so much about my research, provided patient guidance, and responded to my questions so promptly during my Co - Op study at Ford Motor Company. Special thanks to Drs. Tsung - Yu Pan and Allen Roche of Vinci Technology Corporation with whom I have ha d the pleasure of working on the NSF - SBIR project. I must, however, give special thanks to Dr. Pan, who truly went above and beyond in providing me with advice and support on matters both academic and otherwise . Thanks to F. H. LePrevos t and Aaron Morrow of Johnson Manufacturing Company for providing materials and advice for aluminum and solder experimentation . Much appreciation must also go out to the ChEMS faculty that I took classes from, the ChEMS staff who facilitated my research and study need s, including Donna Fernandez, Jennifer Keddle, and Tiffany Owen , and my current lab mates for their support and friendship, including Yuxiang Zhong, Yuelin Wu, and David Vogelsang . vi TABLE OF CONTENTS LIST OF TABLES ................................ ................................ ................................ ....................... ix LIST OF FIGURES ................................ ................................ ................................ ..................... xi KEY TO ABBREVIATION S ................................ ................................ ................................ .. xvii KEY TO SYMBOLS ................................ ................................ ................................ ................ xviii Chapter 1 Introduction ................................ ................................ ................................ ................. 1 1.1 Solidification of Aluminum Silicon Alloys ................................ ................................ .......... 1 1.1.1 Solidific ation of Primary Aluminum Phase ................................ ................................ ... 3 1.1.2 Al - Si Eutectic Solidification ................................ ................................ ........................... 4 1.2 Effects of grain refinement and chemical modification ................................ ........................ 6 1.2.1 Strengthening in Al - Si Alloys ................................ ................................ ......................... 6 1.2.2 Ductility Improvement in Al - Si Alloys ................................ ................................ ......... 13 1.3 Development of Cast Aluminum Alloys Using Nano - structured Chemicals ..................... 19 1.3.1 Overall Goal ................................ ................................ ................................ ................ 19 1.3.2 Motivation ................................ ................................ ................................ .................... 19 1.3.3 Fundamental Idea and Specific Aims ................................ ................................ .......... 26 Chapter 2 Sn - Cu Binary Alloys Refinement ................................ ................................ ............ 28 2.1 Background ................................ ................................ ................................ ......................... 28 2.2 Experimental ................................ ................................ ................................ ....................... 29 2.2.1 Preparation of Sn/Cu Solder Wires ................................ ................................ ............. 29 2.2.2 Coil Creep Testing ................................ ................................ ................................ ....... 30 2.2.3 Polarized Optical Microscopy ................................ ................................ ..................... 30 2.2.4 Differential Scanning Calorimetry ................................ ................................ .............. 31 2.2.5 Indentation Microhardness ................................ ................................ .......................... 31 2.2.6 Orientation Imaging Microscopy ................................ ................................ ................. 31 2.3 Results and Discussions ................................ ................................ ................................ ...... 32 2.3.1 Creep Behavior ................................ ................................ ................................ ............ 32 2.3.2 Microstructure ................................ ................................ ................................ ............. 34 2.3.3 Solidification of Eutectic Constituent ................................ ................................ .......... 35 2.3.4 Indentation Microhardness ................................ ................................ .......................... 37 2.3.5 EBSD Results on Tin - rich Grains ................................ ................................ ................ 38 2.3.6 Texture of Tin - ri ch Grains ................................ ................................ ........................... 40 2.4 Conclusions ................................ ................................ ................................ ......................... 41 Chapter 3 Sn - based Ternary Alloys ................................ ................................ .......................... 43 3.1 Background ................................ ................................ ................................ ......................... 43 3.2 Experimental ................................ ................................ ................................ ....................... 44 3.2.1 Preparation of Solid Wires ................................ ................................ .......................... 44 3.2.2 C oil Creep Testing ................................ ................................ ................................ ....... 44 3.2.3 Polarized Optical Microscopy ................................ ................................ ..................... 45 vii 3.3 Results and Discussions ................................ ................................ ................................ ...... 45 3.3.1 Pure Tin ................................ ................................ ................................ ....................... 45 3.3.2 Sn - Ag - Cu Alloys ................................ ................................ ................................ ........... 47 3.3.3 Sn - Cu - Sb Alloys ................................ ................................ ................................ ........... 50 3.3.4 Sn - Cu - Zn Alloys ................................ ................................ ................................ ........... 53 3.4 Conclusions ................................ ................................ ................................ ......................... 56 Chapter 4 Al - 12Si Alloys with TSP Treatment ................................ ................................ ........ 57 4.1 Background ................................ ................................ ................................ ......................... 57 4.2 Experimental ................................ ................................ ................................ ....................... 58 4.2.1 Materials ................................ ................................ ................................ ...................... 58 4.2.2 Surface Treatment of A4047 powders ................................ ................................ .......... 58 4.2.3 Preparation of A4047 Ingots ................................ ................................ ....................... 58 4.2.4 Microscopy ................................ ................................ ................................ ................... 60 4.2.5 Tensile Test ................................ ................................ ................................ .................. 60 4 .2.6 Indentation Microhardness ................................ ................................ .......................... 60 4.2.7 Thermal Analysis ................................ ................................ ................................ ......... 60 4.3 Results and Discussion ................................ ................................ ................................ ....... 61 4.3.1 Mechanical Properties ................................ ................................ ................................ . 61 4.3.2 As - cast Microstructure ................................ ................................ ................................ . 63 4.3.3 Fracture Surface ................................ ................................ ................................ .......... 67 4.3.4 Solidification ................................ ................................ ................................ ................ 70 4.3.5 Influence of Treatment Time ................................ ................................ ........................ 74 4.3.6 Sequential Dilution ................................ ................................ ................................ ...... 77 4.3.7 Proposed Modification Mechanism ................................ ................................ ............. 83 4.4 Conclusions ................................ ................................ ................................ ......................... 85 Chapter 5 Development of Al - Si - TSP Master All oys ................................ .............................. 86 5.1 Background ................................ ................................ ................................ ......................... 86 5.2 Experimental ................................ ................................ ................................ ....................... 87 5.2.1 Materials and Melt Preparation ................................ ................................ .................. 87 5.2.2 TSP Incorporation P rocess ................................ ................................ .......................... 87 5.2.3 Casting ................................ ................................ ................................ ......................... 87 5.2.4 Characteristics ................................ ................................ ................................ ............. 88 5.3 Results and Discussion ................................ ................................ ................................ ....... 88 5.3.1 Melt Characterization ................................ ................................ ................................ .. 88 5.3.2 Form of Master ................................ ................................ ................................ ............ 89 5.3.3 TSP Concentration in Master ................................ ................................ ...................... 95 5.3.4 Amount of Master ................................ ................................ ................................ ....... 100 5.5 Conclusions ................................ ................................ ................................ ....................... 103 Chapter 6 Aluminum Silicon Binary Alloys Modification ................................ .................... 104 6.1 Background ................................ ................................ ................................ ....................... 104 6.2 Experimental ................................ ................................ ................................ ..................... 105 6.2.1 Materials ................................ ................................ ................................ .................... 105 6.2.2 TSP Master Alloy Preparation ................................ ................................ ................... 105 6.2 .3 Casting ................................ ................................ ................................ ....................... 105 viii 6.2.4 Characteristics ................................ ................................ ................................ ........... 107 6.3 Results and Discussion ................................ ................................ ................................ ..... 108 6.3.1 Melt Characterization ................................ ................................ ................................ 108 6.3.2 Solidification and Melting ................................ ................................ .......................... 109 6.3.3 As - cast Microstructures ................................ ................................ ............................. 113 6.3.4 Tensile Properties ................................ ................................ ................................ ...... 115 6.3.5 Additive Stability ................................ ................................ ................................ ........ 1 17 6.4 Conclusions ................................ ................................ ................................ ....................... 125 Chapter 7 Commercial Cast Aluminum Silicon Based Alloys Modification ...................... 126 7.1 Background ................................ ................................ ................................ ....................... 126 7.2 Experimental ................................ ................................ ................................ ..................... 127 7.2.1 Materials ................................ ................................ ................................ .................... 127 7 .2.2 TSP Master Alloy Preparation ................................ ................................ ................... 127 7.2.3 Casting ................................ ................................ ................................ ....................... 127 7.2.4 Characteristics ................................ ................................ ................................ ........... 128 7.3 Results and Discussion ................................ ................................ ................................ ..... 130 7.3.1 Melting Characterization ................................ ................................ ........................... 130 7.3.2 Casting Properties ................................ ................................ ................................ ..... 132 7.3.3 Solidification and Melting ................................ ................................ .......................... 135 7.3.4 As - cast Microstructures ................................ ................................ ............................. 144 7.3.5 Tensile Properties ................................ ................................ ................................ ...... 148 7.3.6 High Cycle Fatigue ................................ ................................ ................................ .... 151 7.4 Conclusions ................................ ................................ ................................ ....................... 153 Chapter 8 Summary a nd Future Work ................................ ................................ .................. 154 8.1 Summary ................................ ................................ ................................ ........................... 154 8.2 Future Work ................................ ................................ ................................ ...................... 155 8.2.1 Understanding modification mechanism wi th TSP addition. ................................ .... 155 8.2.2 Effect of TSP on the porosity distribution and fatigue in cast aluminum alloys ........ 155 8.2.3 Application of TSP in foundries ................................ ................................ ................. 156 APPENDIX ................................ ................................ ................................ ................................ 157 BIBLIOGRAPHY ................................ ................................ ................................ ..................... 172 ix LIST OF TABLES Table 2 - 1 Comparison of Solidific ation Temperature, Melting Temperature, and Undercooling for Sn - Cu Wires. ................................ ................................ ................................ ........................... 35 Table 3 - 1 Chemical Composition in wt% of All Wires. ................................ ............................... 44 Table 3 - 2 Stress Exponent and Activation Energy of Pure Sn and TSP - treated Sn Wires. ......... 46 Table 3 - 3 Stress Exponent and Activation Energy of Sn - 4.7Ag - 1.7Cu and TSP - treated Sn - 4.7Ag - 1.7Cu Wires. ................................ ................................ ................................ ................................ . 49 Table 3 - 4 Stress Exponent and Activation Energy of Sn - 3.5Cu - 7.5Sb and TSP - treated Sn - 3.5Cu - 7.5Sb Wires. ................................ ................................ ................................ ................................ .. 52 Table 3 - 5 Stress Exponent and Activatio n Energy of Sn - 15Cu - 3Zn and TSP - treated Sn - 15Cu - 3Zn Wires. ................................ ................................ ................................ ................................ ..... 55 Table 4 - 1 Overall Composition of TSP in the A4047 Ingots. ................................ ...................... 59 Table 4 - 2 Eutectic Reaction Temperatures for A4047 with and without TSP Extracted from Cooling Curves. ................................ ................................ ................................ ............................ 73 Table 4 - 3 Microhardness of the TSP - treated A4047 Alloys. ................................ ....................... 76 Table 4 - 4 Microhardness of the Diluted A4047 Alloys. ................................ ............................... 78 Table 4 - 5 Eutectic Reaction Temperatures Extracted from Cooling Curves. .............................. 82 Table 5 - 1 Chemical Composition in wt% of the Al - 7.5Si Base Alloy. ................................ ........ 88 Table 5 - 2 Naming Conventions of All Cast Specimens for Master Form Study. ........................ 89 Table 5 - 3 SDAS of the Al - 12Si Alloys with and without TSP. ................................ ................... 92 Table 5 - 4 Naming Conventions of All Cast Specimens for TSP Concentration Study. ............... 95 Table 5 - 5 Naming Conventions of All cast Specimens for Master Content Study. ................... 100 Table 6 - 1 Naming Conventions of All C ast Specimens. ................................ ............................ 106 Table 6 - 2 Chemical Composition in wt% of the Al - 7.5Si Alloy with and without TSP. ........... 108 Table 6 - 3 Selected Characteristic Temperatures Taken from Cooling Curves. ......................... 111 Table 6 - 4 Selected Characteristic Temperatures Taken from DSC Heating Curves. ................. 112 x Table 7 - 1 Chemi cal Composition in wt% of the Aural TM 2 Alloys. ................................ ........... 131 Table 7 - 2 Chemical Composition in wt% of the W319 Alloys. ................................ ................. 131 Table 7 - 3 Selected Characteristic Temperatures Taken from Cooling Curves. ......................... 140 Table 7 - 4 High Cycle Fatigue Stress of Heat - treated Aural TM 2 Alloys at Room Temperature. 152 xi LIST OF FIGURES Figure 1 - 1 Aluminum Silicon Binary Phase Diagram [10]. ................................ ........................... 2 Figure 1 - 2 Type of Al - Si Eutectic Structures as a Function of Growth Rate (V) and Temperature Gradient (G): ................................ ................................ ................................ ................................ ... 5 Figure 1 - 3 The Aluminum Rich - side of the Al - Ti Phase Diagram [29]. ................................ ........ 8 Figure 1 - 4 Al 3 Ti Nucleates on Aluminum Grain [30]. ................................ ................................ ... 9 Figure 1 - 5 The Aluminum Rich - side of the Al - B Phase Diagram [31]. ................................ ....... 12 Figure 1 - 6 Chemical Structure of (a) Polyhedral Oligom eric Silsesquioxane (POSS) and (b) Trisilanol POSS. ................................ ................................ ................................ ........................... 21 Figure 2 - 1 Creep Experimental Setup: (a) Stereo Micrographs of Wire Winding and (b) Schematic of Coil Creep Setup. ................................ ................................ ................................ .... 30 Figure 2 - 2 Stretch Rates at Various Temperatures: (a) Sn - 0.7Cu Wire without and with T SP Additions and (b) Sn - 3Cu Wire without and with TSP Additions. ................................ .............. 33 Figure 2 - 3 Polarized Optical Micrographs of Sn - Cu Solder Wires: (a) Sn - 0.7Cu Wire, (b) TSP - treated Sn - 0.7Cu Wire, (c) Sn - 3Cu Wire, and (d) TSP - treated Sn - 3Cu Wire. ............................ 36 Figure 2 - 4 Knoop Hardness Values of Sn - Cu Wires. ................................ ................................ ... 37 Figure 2 - 5 C - axis orientation maps of Sn - Cu solder wires: (a) Sn - 0.7Cu wire, (b) TSP - treated Sn - 3Cu wire, (c) Sn - 3Cu wire, and (d) TSP - treated Sn - 3Cu wire ................................ ................ 39 Figure 2 - 6 (001) and (100) discrete and density pole Fig.s of Sn - Cu solder wires: (a) Sn - 0.7Cu wire, (b) TSP - treated Sn - 3Cu wire, (c) Sn - 3Cu wire, and (d) TSP - treated Sn - 3Cu wire ............. 41 Figure 3 - 1 Stretch Rates of Pure Sn and TS P - treated Sn Wires at Various Temperatures. ......... 46 Figure 3 - 2 Optical Micrographs of Pure Tin Solder Wires: (a) Pure Sn Wire and (b) TSP - tr eated Sn Wire. ................................ ................................ ................................ ................................ ........ 47 Figure 3 - 3 Stretch Rates of Sn - 4.7Ag - 1.7Cu and TSP - treated Sn - 4.7Ag - 1.7Cu Wires at Various Temperatures. ................................ ................................ ................................ ................................ 48 Figure 3 - 4 Optical Micrographs of Sn - 4.7Ag - 1.7Cu Solder Wires: (a) Sn - 4.7Ag - 1.7Cu Wire and (b) TSP - treated Sn - 4.7Ag - 1.7Cu Wire. ................................ ................................ ........................ 50 Figure 3 - 5 Stretch Rates of Sn - 3.5Cu - 7.5Sb and TSP - treated Sn - 3.5Cu - 7.5Sb Wires at Various Temperatures ................................ ................................ ................................ ................................ . 51 xii Figure 3 - 6 Optical Micrographs of Sn - 3.5Cu - 7.5Sb Solder Wires: (a) Sn - 3.5Cu - 7.5Sb Wire and (b) TSP - treated Sn - 3.5Cu - 7.5Sb Wire. ................................ ................................ ......................... 53 Figure 3 - 7 Stretch Rates of Sn - 15Cu - 3Zn and TSP - treated Sn - 15Cu - 3Zn Wires at Various Temperatures. ................................ ................................ ................................ ................................ 54 Figure 3 - 8 Optical Micrographs of Sn - 15Cu - 3Zn Solder Wires: (a) Sn - 15Cu - 3Zn Wire and (b) TSP - treated Sn - 15Cu - 3Zn Wire. ................................ ................................ ................................ .. 56 Figure 4 - 1 Ultimate Tensile Strength (a) and Elongation to Failure (b) of Ingots Made of A4047 Powders without and with TSP Treatment. ................................ ................................ .................. 62 Figure 4 - 2 Optical Micrographs of Al - Si Alloys Showing the Effect of TSP on the Primary Al and Eutectic Si Morphology of the as - Cast Ingot: (a) The A4047 Contro l with an Objective Lens of 5x, (b) The TSP - treated A4047 with an Objective Lens of 5x, (c) The A4047 Control with an Objective Lens of 20x, (d) The TSP - treated A4047 with an Objective Lens of 20x, (e) The A4047 Control with an Objective Lens of 100x, and (f) The TSP - treated A4047 with an Objective Lens of 100x. ................................ ................................ ................................ ................ 64 Figure 4 - 3 Effect of TSP Treatment on the Secondary Dendrite Arm Spacing (SDAS) of the A4047 Ingots. ................................ ................................ ................................ ................................ 65 Figure 4 - 4 ImageJ Analysis of (a) the A4047 Ingot and (b) the TSP - treated A4047 Ingot. ........ 66 Figure 4 - 5 A Fractured A4047 Tensile Bar: (a) Photograph of Fractured Sample, (b) Stereo Photograph of Fracture Surface, and (c) Back - scattered Electron Micrograph of Fracture Surface. ................................ ................................ ................................ ................................ ....................... 68 Figure 4 - 6 A Fractured TSP - treated A4047 Tensile Bar: (a) Photograph of Frac tured Sample, (b) Stereo Photograph of Fracture Surface, and (c) Back - scattered Electron Micrograph of Fracture Surface. ................................ ................................ ................................ ................................ ......... 69 Figure 4 - 7 C ooling Curves of A4047 and TSP - treated A4047 Alloys. ................................ ........ 72 Figure 4 - 8 Optical Micrograph of an Ingot Made of the A4047 Powders: (a) without TSP Treatment, (b) 4 - hour Treatment, (c) 8 - hour Treatment, (d) 24 - hour Treatment, and (e) 9 - day Treatment. ................................ ................................ ................................ ................................ ..... 75 Figure 4 - 9 Optical Micrographs of Diluted A4047 Ingots: (a) 50% Untreated A4047 + 50% Untreated A4047, (b) 50% Untreated A4047 + 50% TSP - treated A4047, (c) 75% Untreated A4047 + 25% TSP - treated A4047, (d) 94% Untreated A4047 + 6% TSP - treated A4047. .......... 80 Figure 4 - 10 Cooling Curves of Diluted TSP - treated A4047 ingots. ................................ ............ 82 Figure 4 - 11 Schematic Illustrating 2D View of the Twin Plane Re - entrant Edge Growth Mechanism. ................................ ................................ ................................ ................................ ... 85 xiii Figure 5 - 1 Optical Micrographs of As - cast Al - 12Si Alloys with and without TSP Using Cold - pressed Puck Process: (a) the Al - 12Si - P Base Alloy with the Objective Lens of 20x, (b) the Al - 12Si - TSP - P with the Objective Lens of 20x, (c) the Al - 12Si - P Base Alloy w ith the Objective Lens of 100x, and (d) the Al - 12Si - TSP - P with the Objective Lens of 100x. ............................... 90 Figure 5 - 2 Optical Micrographs of As - cast Al - 12Si Alloys with and without TSP using Pre - melted Ingot Process: (a) the Al - 12Si - I Base Alloy with the Objective Lens of 20x, (b) the Al - 12Si - TSP - I with the Objective Lens of 200, (c) the Al - 12Si - I Base Alloy with the Objective Lens of 100x, and (d) the Al - 12Si - TSP - I with the Objective Lens of 100x. ................................ ......... 91 Figure 5 - 3 Optical Micrographs of As - cast Al - 7.5Si Alloys with and without TSP Using Cold - pressed Puck Process: (a) the Al - 7.5Si - P Base Alloy with the Objective Lens of 20x, (b) the Al - 7.5Si - TSP - P with the Objective Lens of 20x, (c) the Al - 7.5Si - P Base Alloy with the Objective Lens of 100x, and (d) the Al - 7.5Si - TSP - P with the Objecti ve Lens of 100x. .............................. 93 Figure 5 - 4 Optical Micrographs of As - cast Al - 7.5Si Alloys with and without TSP Using Pre - melted Ingot Process: (a) t he Al - 7.5Si - I Base Alloy with the Objective Lens of 20x, (b) the Al - 7.5Si - TSP - I with the Objective Lens of 200x, (c) the Al - 7.5Si - I Base Alloy with the Objective Lens of 100x, and (d) the Al - 7.5Si - TSP - I with the Objective Lens of 100x. ............................... 94 Figure 5 - 5 Optical Micrographs of the A4047 - TSP Master with Different Amount of TSP: (a) the A4047 - 2TSP with the Objective Lens of 20x, (b) the A4047 - 2T SP with the Objective Lens of 100x, (c) the A4047 - 4TSP with the Objective Lens of 20x, (d) the A4047 - 4TSP with the Objective Lens of 100x, (e) the A4047 - 6TSP with the Objective Lens of 20x, and (f) the A4047 - 6TSP with the Objective Lens of 100x. ................................ ................................ ........................ 97 Figure 5 - 6 Optical Micrographs of As - cast Al - 7.5Si Alloys with and without TSP: (a) the Al - 7.5Si Base Alloy with the Objective Lens of 20x, (b ) the Al - 7.5Si - TSP - i with the Objective Lens of 20x, (c) the Al - 7.5Si Base Alloy with the Objective Lens of 100x, and (d) the Al - 7.5Si - TSP - i with the Objective Lens of 100x. ................................ ................................ ................................ .. 98 Figure 5 - 7 Optical Micrographs of As - cast Al - 7.5Si Alloys with TSP: (a) the Al - 7.5Si - TSP - ii with the Objective Lens of 20x, (b) the Al - 7.5Si - TSP - iii with the Objective Lens of 20x, (c) the Al - 7.5Si - TSP - ii with the Objective Lens of 100x, and (d) the Al - 7.5Si - TSP - iii with the Objective Lens of 100x. ................................ ................................ ................................ ................................ 99 Figure 5 - 8 Optical Mic rographs of As - cast Al - 7.5Si alloys with TSP: (a) Al - 7.5Si with the Objective Lens of 20x, (b)Al - 7.5Si - TSP - 25 with the Objective Lens of 20x, (c) Al - 7.5Si with the Objective Lens of 100x, (d) Al - 7.5Si - TSP - 25 with the Objective Lens of 100x. ...................... 101 Figure 5 - 9 Optical Micrographs of As - cast Al - 7.5Si Alloys with TSP: (a) Al - 7.5Si - TSP - 10 with the Objective Lens of 20x, (b) Al - 7.5Si - TSP - 5 with th e Objective Lens of 20x, (c) Al - 7.5Si - TSP - 10 with the Objective Lens of 100x, (d) Al - 7.5Si - 6TSP - 5 with the Objective Lens of 100x. .. 102 Figure 6 - 1 Cooling Curve and its First Derivative of (a) Al - 7.5Si and (b) Al - 7.5Si - TSP - 1.5h. 110 Figure 6 - 2 DSC Heating Curves of the A l - 7.5Si Alloy with and without TSP. ......................... 112 xiv Figure 6 - 3 Optical Micrographs of the Al - 7.5Si Alloy with and without TSP: (a) Al - 7.5Si, (b) Al - 7.5Si with TSP Additions after 1.5 - hour Hold in 720 o C Furnace. ................................ ....... 114 Figure 6 - 4 Stress - strain Curves of the Al - 7.5Si Alloy with and without TSP at Room Temperature. ................................ ................................ ................................ ............................... 115 Figure 6 - 5 Tensi le Strength (a) and Ductility (b) of the Al - 7.5Si Alloy with and without TSP. 116 Figure 6 - 6 Cooling Curves of the Al - 7.5Si, Al - 7.5Si - TS P - 1.5h, and Al - 7.5Si - TSP - 4h Alloys. 118 Figure 6 - 7 DSC Heating Curves of the Al - 7.5Si, Al - 7.5Si - TSP - 1.5h, and Al - 7.5Si - TSP - 4h Alloys. ................................ ................................ ................................ ................................ ......... 119 Figure 6 - 8 Optical Micrographs of the Al - 7.5Si and Al - 7.5Si - TSP - 4h Alloys. (a) Al - 7.5Si, (b) Al - 7.5Si With T SP Additions After 4 - hour Hold in 720 o C Furnace. ................................ ........ 120 Figure 6 - 9 Tensile Strength (a) and Ductility (b) of the Al - 7.5Si, Al - 7.5Si - TSP - 1. 5h, and Al - 7.5Si - TSP - 4h Alloys. ................................ ................................ ................................ .................. 121 Figure 6 - 10 Cooling Curves of the Al - 7.5Si, Al - 7.5Si - TSP - 4h, and Al - 7.5Si - TSP - 4h - R - X. X=24, 48, 72, and 192h. ................................ ................................ ................................ ......................... 123 Figure 6 - 11 Optical Micrographs of the Al - 7.5Si, Al - 7.5Si - TSP - 4h, and Al - 7.5Si - TSP - 4h - R - X. X=24, 48, 72, and 192h: (a) Al - 7.5Si, (b) Al - 7.5Si with TSP Additions after 4 - hour Hold in 720 o C Furnace, (c) Al - 7.5Si with TSP Additions after Remelting 4 - hour - hold Al - 7.5S i - TSP Alloy and Additional 24 - hour Hold in 720 o C Furnace, (d) Al - 7.5Si with TSP Additions after Remelting 4 - hour - hold Al - 7.5Si - TSP Alloy and Additional 48 - hour Hold in 720 o C Furnace, (e) Al - 7.5Si with TSP Additions after Remelting 4 - hour - hold Al - 7.5Si - TS P Alloy and Additional 72 - hour Hold in 720 o C Furnace, (f) Al - 7.5Si with TSP Additions after Remelting 4 - hour - hold Al - 7.5Si - TSP Alloy and Additional 192 - hour Hold in 720 o C Furnace. ................................ ... 124 Figure 7 - 1 A Top - view Image of Aural TM 2 Castings: Aural TM 2 Control, Aural TM 2 TSP, and Aural TM 2 Sr Alloys. ................................ ................................ ................................ ................ 133 Figure 7 - 2 A Top - view Image of W319 Castings: W319 Control and W319 - TSP Alloys. ....... 134 Figure 7 - 3 Typical DSC Endothermic (Top) and Exothermic (Bottom) Curves of Aural TM 2, Containing the Primary Al and Al - Si Eutectic Microconstituent. ................................ .............. 136 Figure 7 - 4 DSC Results of Aural TM 2 (Blue), Aural TM 2 - TSP(Red), and Aural TM 2 - Sr (Green) during Heating (Top) and Cooling (Bottom) Cycles. ................................ ................................ . 137 Figure 7 - 5 Cooling Curve and its First Derivative of (a) W319 and (b) W319 - TSP. ................ 13 9 Figure 7 - 6 Typical DSC Endothermic Curve of W319. Al: The Primary Al, Si: Al - Si E : Al 2 Cu, Q: Al 4 Cu 2 Mg 8 Si 7 and - Fe: Al 15 (Fe, Mn) 3 Si 6. ................................ ................................ 142 xv Figure 7 - 7 DSC Results of W319 (Blue), W319 - TSP(Red), and W319 - Sr (Green) during Heating Cycle. ................................ ................................ ................................ ................................ .......... 143 Figure 7 - 8 Optical Micrographs of the As - cast Condition: (a) Aural TM 2 Base, (b) Aural TM 2 with TSP, and (c) Aural TM 2 with Sr. ................................ ................................ ................................ .. 144 Figure 7 - 9 SEM Micrographs of the Deep - etched Aural TM 2 Alloys: (a) Aural TM 2 Base, (b) Aural TM 2 with TSP, and (c) Aural TM 2 with Sr. ................................ ................................ ........ 145 Figure 7 - 10 EBSD Micrographs of the As - cast Aural TM 2 Alloys: (a) Aural TM 2 Base, (b) Aural TM 2 with TSP, and (c) Aural TM 2 with Sr. ................................ ................................ ........ 145 Figure 7 - 11 Optical Micrographs of the W319 Alloy with and without TSP: (a) W319 and (b) W319 with TSP Additions. ................................ ................................ ................................ ......... 147 Figure 7 - 12 Stress - strain Curves of the Aural TM 2 Alloy with and without TSP at (a) Room Temperature and (b) 300 o C. ................................ ................................ ................................ ........ 149 Figure 7 - 13 Yield Strength, Ultimate Tensile Strength, and Ductility of the Aural TM 2 Alloy with and without TSP Additions at Room Temperature and 3 00 o C. ................................ .................. 150 Figure 7 - 14 S - N Curves of Heat - treated Aural TM 2 Alloys: (a) Aural TM 2 without Sr, (b) Aural TM 2 with TSP, and (c) Aural TM 2 with Sr. ................................ ................................ ....................... 151 Figure 7 - 15 Fracture Surface of Fatigue Specimen: (a) Aural 2 without Sr, (b) Aural 2 with T SP, and (c) Aural 2 with Sr Alloys. ................................ ................................ ................................ ... 152 Figure A - 1 Cartoon of a Four - point Flexure Fixture. ................................ ................................ . 159 Figure A - 2 A Four - point Bend Fixture with a Bended Specimen: (a) Side - view and (b) Top - view. ................................ ................................ ................................ ................................ ............ 160 Figure A - 3 Macro of a Four - point Bend specimen: (a) A4047 and (b) A4047 - TSP. ................. 162 Figure A - 4 Opical Micrographs of As - cast A4047 Alloys with and without TSP Treatment Using Dip - coating Process: (a) the A4047 Base Alloy with the Objective Lens o f 20x, (b) the T4047 - TSP with the Objective Lens of 20x, (c) the A4047 Base Alloy with the Objective Lens of 100x, and (d) the A4047 - TSP with the Objective Lens of 100x. ................................ ......................... 163 Figure A - 5 Stereo Micrographs of Bended A4047 Alloys with and without TSP Treatment Using Dip - coating Process: (a) the A4047 Base Alloy with the Objective Lens of 7x, (b) the A4047 - TSP with the Objective Lens of 7x, (c ) the A4047 Base Alloy with the Objective Lens of 10x, and (d) the A4047 - TSP with the Objective Lens of 10x. ................................ ........................... 164 Figure A - 6 Secondary ( Left) and Backscattered (Right) Electron Images of the A4047 Ingots. A 1.9mm - long Crack was Formed from the Notch. ................................ ................................ ....... 166 xvi Figure A - 7 Secon dary (Left) and Backscattered (Right) Electron Images of the TSP - treated A4047 Ingots. A 0.78mm - long crack was formed from the notch. ................................ ............ 167 Figure A - 8 Strains Induced during Bending of a Beam. ................................ ............................ 169 Figure A - 9 Stereo Micrographs of (a) the A4047 Ingot and (b) the TSP - treated A4047 Ingot Analyzed by ImageJ. ................................ ................................ ................................ ................... 170 xvii KEY TO ABBREVIATIONS Al Aluminum Cu Copper Mg Magnesium Si Silicon Sn Tin Sr Strontium FCC Face - Centered C ubic OM Optical Microscope SEM Scanning Electron Microscopy VH Vickers Hardness UTS Ultimate Tensile Strength POSS Polyhedral O ligomeric S ilse squioxane TSP TriSilanol POSS IMCs Intermetallic Compounds SDAS Secondary Dendrite Arm Spacing xviii KEY TO SYMBOLS a Lattice parameter r* the critical radius for homogeneous nucleation the latent heat of fusion per unit volume of solid formation the solid/liquid interfacial free energy the equilibrium solidification temperature the undercooling V growth rate G temperature gradient S slope of the liquidus line y Yield strength Strain d the average grain diameter the critical nucleus radius for heterogeneous nucleation the free energy barrier the number of clusters with radius r the free energy change per unit volume of solid formation the solid/liquid interfacial free energy wetting angle cluster free energy total number of atoms k Boltzmans constant xix T Temperature interparticle spacing the lamellar spacing at the extremum undercooling point the ratio of mean to extremum spacing T Al nuc the primary Al nucleation temperature T Al min the primary Al temperature minimum (dT/dt=0). T Al max the primar y Al temperature maximum (dT/dt=0). T Al - Si nuc the Al - Si eutectic nucleation temperature. T Al - Si arr the Al - Si eutectic arrest temperature. T sol nuc the Cu IMCs nucleation temperature. 1 Chapter 1 Introduction Aluminum based casting alloys have been widely used in areas such as automotive, military, aerospace, and alternative energy industries , due to their light - weight and low life - cycle costs characteristics [ 1 - 4 ] . To meet future fuel standards and safety requirements , i t is expected the demand for aluminum alloys in transportation applications will increase dramatically , includ ing powertrain and structural comp onents that were traditionally cast from iron. As compared with the eutectic transformation temperature of Al - Cu (548 °C ) [ 5 ] and Al - Mg(450 °C ) [ 6 ] binary alloys, Al - Si alloys have a higher eutectic t emperature , 577 °C [ 7 ] . Therefore , Al - Si alloys exhibit n arrower solidification ranges , and offer greater fluidity [ 8 , 9 ] , which precede other aluminum casting alloys. The mechanical properties of Al - Si casting alloys, such as ductility and strength, are determined by the metallurgy of solidification. This chapter addresses the role the solidification play s in the Al - Si alloy development and discusses intrinsic effects of alloying elements on solidification and mechanical properties. 1 .1 Solidification of Aluminum Silicon Alloys The evolution of Al - Si alloy microstructure during solidification under both equilibrium and nonequ ilibrium conditions can be predicted from the phase diagram. Al atoms are arranged in a face - centered cubic structure with a lattice spacing of 0.40 nm , while Si crystallizes in a diamond cubic crystal structure wit h a lattice spacing of 0.54 nm [ 7 ] . The difference in these crystal structures results in three equilibrium states in terms of the minimum free energy at a given temperature: primary aluminum phase, Si phase, and Al - Si eutectic microconstituent . According to free energy curves of Al - Si alloys obtained at different temperatures, the equilibrium phase diagram is generated as shown in Figure 1 - 1 [ 10 ] . 2 Figure 1 - 1 Aluminum Silicon Binary Phase Diagram [10] . 3 In Figure 1 - 1 , Al - Si alloys containing up to 12.6 wt . % Si and over 12.6 wt . % Si are referred to as hypoeutectic alloys and hypereutectic alloys, respectively. Hypoeutectic Al - Si alloys are more suitable for casting than the hypereutectic alloys, since the larger size and non - uniform distribution of the brittle primary Si phase in hypereutectic alloys result in very poor ductility. Typically, solidification of hypoeutectic Al - Si alloys consists of solidification of the primary Al phase followed by the Al - Si eutectic re action. 1.1.1 Solidification of Primary Aluminum Phase The solidification of primary aluminum occurs by two sequen tial microstructure evolution s : nucleation and formation of a dendritic network [ 11 ] . Nucleation is the formation of nuclei from molten material. Based on the free energy change from liquid state to solid state , classical nucleation theory was developed in the early twenty century [ 12 - 18 ] . There are two types of nucleati on: homogeneous and heterogeneous. Homogeneous nucleation typically occur s when the undercooling become s large enough to ca us e the formation of a stable nucleus [ 19 ] . The size of the critical radius r* for ho mogen e ous nucleation [ 20 ] is given by Equation 1 - 1 where is the latent heat of fusion per unit volume of solid formation, is the solid/liquid interfacial free energy, is the equilibrium solidification temperature, and is the undercooling when the liquid temperature is T. Heterogeneous nucleation occurs when there are impurities, either suspended in the liquid or on the walls of the container that holds the liquid, providing a surface on which the solid can form. Therefor e, m uch less undercooling is required to achieve the critical size in heterogeneous 4 nucleation. Homogeneous nucleation never happens in practice , while h eterogeneous nucleation typically occurs for the primary aluminum phase [ 20 ] . Once the solid nuclei form s , growth starts as more atoms attach to the solid surface until fully solidif y . The growing process generates heat due to the reduced ent halpy of the aluminum from solid solution to the melt [ 21 , 22 ] , which creates a negative thermal gradient in front of the solid/liquid interface . In the meantime, Si is partition ed into the liquid ahead of the solid/liquid interface as its solubility in the s olid solution is much less than that in liquid, resulting in a corresponding variation in the liquidus temperature based on Figure 1 - 1 . As a result of the supercooled liquid ahead of interface, a small solid protuberance forms [ 23 ] . The temperature drop from interface to protuberance drives the tip to grow, hence form ing the primary dendrite arms. The Si partition perpendicular to the primary arms allows more protuberances to from along the primary arms, which leads to the formation of secondary dendrite arms. The fast growth directions are generally for aluminum , because of the anisotropic surface tension [ 24 ] . As solidification proceeds, dendritic tips become sharper so as to minimize the solid/liquid interfacial energy. 1 .1. 2 Al - Si Eutectic Solidification T he formation of the - phase increases the Si level in the remaining liquid. The eutectic reaction takes place when the Si level in the liquid reaches 12.6 wt.%. Two solid pha ses, the Al - rich and the Si phase, form cooperatively as an anomalous structure fr om liquid during eutectic reaction [ 25 ] . Day and Hellawell [ 26 ] proposed the morphology and crystallography of Si in Al - Si alloys can be classified into regions A, B, and C , as shown in Figure 1 - 2 , based on the growth rate (V) and temperature gr adient (G). 5 Figure 1 - 2 Type of Al - Si E utectic S tructures as a F unction of G rowth R ate (V) and T emperature G radient (G) : ( i ) R egion A (V=0.28 µ m/s and G=30 °C /mm); ( ii ) R egion B (V=0.28 µ m/s and G=20 °C /mm); ( iii ) R egion B (V=0.28 µ m/s and G=3 °C /mm); ( iv ) R egion C (V=3.7 µ m/s and G=3 °C /mm); ( v ) R egion C (water quenched ) [ 26 ] . 6 In Figure 1 - 2 ( i ), massive crystals of planar aluminum and faceted Si phas es are growing by long - range diffusion in region A . The fiber - like Si with round cross - section typically occurs when the temperature gradient is less than 20 °C /mm (Figure 1 - 2 ( ii ) ). However, the morphology of Si in region B is sensitive to G, and the fibrous Si forms lateral branches at the lower temperature gradient (Figure 1 - 2 ( iii ) ). I rregular Si crystals form in region C when the growth rate and temperatu re gra dients fall into a limited range of value (Figure 1 - 2 ( iv ) ). Depending on the solidification conditions, finer structures form at rapid freezing rate of quenched alloy in Figure 1 - 2 ( v ) . Day and Hellawell found that the Si structure grows at a fluctuating duplex solid - liquid front at a rapid freezing rate to form fibrous structure, whereas at relatively low rate s these solid - liquid fronts are irregular and non - isothermal to form flaky structure. Nevertheless, t he solidifica tion condition in region C is similar to that commonly used in laboratory and industrial practice [ 26 ] . Therefore, the microstructure associated with region C is selected as target to study effects on solidification and mechanical properties. 1.2 Effects of grain refinement and chemical modification T o improve the overall mechanical properties of Al - Si alloys , it is important to understand the development of Al - Si microstructures with desirable properties. There are several approaches to alter the as - solidified microstructure of the alloys, which control the mechanical properties. The most significant of which are grain refinement of primary Al to enhance strength and chemical modification of eutectic Si to improve ductility. 1.2.1 Strengthening in Al - Si Alloys Grain refinement is one of the fundamental strengthening approaches to en hance the hardness and strength of Al - Si alloys [ 27 ] . F ine - grained polycrysta l ha ve higher hardness and strength than coarse grained, because the dislocat ion motion during plastic deformation is 7 impeded by a greater total grain boundary area. The yield strength y depends on grain size according to the Hall - Petch equation [ 28 ] y = 0 +k y d - 1/2 Equation 1 - 2 where d is the average grain diameter, and 0 and k y are material consta nt . Grain refiners were typically inoculated in the melt before casting , resulting in a fine equiaxed microstructure s . The refined microstructure was achieved by inoculation during the formation and growth of primary aluminum phase. Titanium was first used to refin e aluminum by the nucleation on Al 3 Ti, which can be explained from the Al - Ti phase diagram in Figure 1 - 3 [ 29 ] . The system presents a peritectic reaction. Liquid + Al 3 Ti Solid (FCC) Equation 1 - 3 Ba ckerud explained how Ti 3 Al nucleates aluminum grain s, as shown in Figure 1 - 4 [ 30 ] . Al - Ti master alloys with many Al 3 Ti particles were inoculated into the aluminum melt. The micro - sized Al 3 Ti crystal s in the master alloys start to dissolve into the liquid aluminum after a few minutes, as the black dot is shown in Figure 1 - 4 (1) and (2). Once the Ti concentration is enriched at the surface of liquid aluminum, Al 3 Ti crystal solidifies first at hyperperitectic concentration ac cording to the Al - Ti phase diagram . Therefore, the solid aluminum nucleate s on the surface of Al 3 Ti in Figure 1 - 4 (3). The Al 3 Ti nuclei are typically located at the interface of the Al 3 Ti crystals and liquid aluminum. After that the primary aluminum crysta l grows around the interface , as shown in Figure 1 - 4 (4) and (5) , s mall solid protuberances begin to grow (6) and dendrites forms (7) as solidification proceeds. 8 Figure 1 - 3 The Aluminum R ich - side of the Al - Ti P hase D iagram [29]. 9 Figure 1 - 4 Al 3 Ti N ucleates on A luminum G rain [30]. 10 There are some concerns about the dissolution of Al 3 Ti crystals. Al 3 Ti dissolves into aluminum when Ti contents fell bel ow the peritectic composition, which causes the loss of refining ability [ 27 ] . T he disappeared grain refining effect with time is known as fade. To overcome this fading issue, Sigworth and Guzowski [ 31 ] used boron as an aluminum refiner, showing that AlB 2 is much more effective in refining aluminum grain si ze as compared to Al 3 Ti. The refinement is caused by a eutectic reaction from the Al - B phase diagram in Figure 1 - 5 [ 31 ] . The AlB 2 nucleus produc ed much smaller grain than Al 3 Ti . In other word s , AlB 2 has a much greater nucleation potency than Al 3 Ti. However, sludge was formed after inoculating B in to the melt [ 32 ] . B has been found to react with other elements, such as Ti, Sr. The refiner - to - refiner/modifier reactions cause sludge or the loss of refinement , which is referred to as The B poisoning issues exhibited a devastating im p act on the soundness of the final products. To solve the Ti fading and B poisoning issues , a combination of B with Ti was developed. Al - Ti - B has been widely used as an excellent grain refiner in aluminum casting industry. However, there is a lot of debate about t he mechanism of the Al - Ti - B grain refinement. The biggest argument focuse s on the right nuclei in the Al - Ti - B master alloy , which is refer red to as the . Two theories under the nucleant paradigm were proposed , which are the nucleant - particle theory and the phase - diagram theory. The nucl eant - particle theory proposed that aluminum nucleates on (Ti/Al)B 2 crystals [ 33 , 34 ] , while the phase diagram theory argued that aluminum nucleates on Al 3 Ti crystals because of the peritectic rea ction [ 35 ] . However, both theories have limitations. The nucleat - partic le theory cannot explain why Al 3 Ti was found at the center of grains at hyperperitectic concentration s . The phase - diagram theories cannot explain grain refinement at hypo p e ri tectic Al - Ti concentrations, since Al 3 Ti dissolves into aluminum for Ti contents below peritectic concentrations. 11 The nucleant paradigm did not provide comprehensive and consistent theo ries to explain grain refinement mechanisms. In 199 0s, the nucleant paradigm was shifted toward the solute paradigm , suggesting that both the nucleant particles and the segregating elements affe ct the grain - refinement process . Easton and StJohn [ 36 ] found that TiB 2 nucleated 99.97% pure aluminum without the addition of solute Ti, but the refined macrostructure w as still columnar with a few equiaxed grains. By adding 0.02% Ti to aluminum containing a TiB 2 substrate, a transition to an equiaxed microstructure was observed with significantly decreased grain size. They also found the same trend for the addition of Ti B 2 and Si solute on the grain size of 99.97% pure aluminum. Similar results were observed by Johnso n [ 37 , 38 ] . Therefo re, it has been experimentally confirmed that both effective nucleants and solute elem e nts are required to obtain grain refinement. The effective nuclea n ts removed the undercooling of substrate, and the solute causes the formation of a constitutional under cooling zone in front of the solidi/liquid interface to facilitate further nucleation. The solute paradigm can also clarify why Al 3 Ti was found at the center of grains in refine d aluminum grains at hypo p e ri tectic Al - Ti concentrations, which can not be expla ined by the nucleant paradigm . The paradigm shift has been experimentally valid ated , however, there is almost no refinement if the Si content is higher than 3% in Al - Si casting alloy s [ 39 ] . A new gain refiner need s to be developed to effectively refine Al - Si alloys with a large range of Si content. 12 Figure 1 - 5 The Aluminum R ich - side of the Al - B P hase D iagram [31]. 13 1.2.2 Ductility Improvement in Al - Si Alloys T he ductility of hypoeutectic Al - Si binary casting alloys is relatively low [ 40 - 43 ] , as compared to steel. This is ma inly because the presence of irregular eutectic Si crystals in the as - cast condition promotes crack initiation and propagation [ 44 ] . Therefore, the manner in which morphology of eutectic Si crystals develops during solidificati on is extremely important to optimize the overall mechanical properties for Al - Si casting alloys. The morphology of irregular eutectic Si crystals is in flaky form, which results in low ductility. In order to increase the ductility without losing strength , trace elements, including alkaline or alkali earth metals such as Na [ 45 ] , Ca [ 46 , 47 ] , or Sr [ 48 - 55 ] , are intentionally added to modify the mo rphology of the Si crystals from the flake - like structure to a fibrous structure. It is difficult to control Na modifications due to its high vapor pressure and fast fading in molten aluminum [ 56 ] . Ca is less effective as a modifier than Na and Sr, while Ca addition s are usually accompanied by an increase in the dissolved hydrogen levels , which adversely affect s casting properties [ 57 ] . Sr has less fading effect in molten aluminum compared to Na and is more effective than Ca [ 58 ] . Therefore, Sr is the most widely used trace element addition for modifying the morphology of eutectic Si crystals to improve ductility. Considering that this modification through chang es in the Si morphology o nly occurs in eutectic reaction , this section focus es on the Si microstructure evolution in eutectic colonies during solidification. The solidification progress of eutectic Si crystals follows two stages: nucleation and growth. A n understanding of the nucleation and growth mechanisms of eutectic Si crystals during solidification, with and without Sr addition s , is critical to the fundamental knowledge of the modification mechanisms of Si crystals with Sr additions and how they relate to the improvement of ductility. 1) Nucleation Kinetics of Eutectic Si Cryst als 14 It is generally accepted that the initial stages of eutectic reactions are controlled by heterogeneous nucleation [ 25 , 59 - 61 ] . According to c lassical nucleation theory [ 20 ] , the critical nucleus radius ( ), the free energy barrier ( ), and the number of clusters with radius r ( ) are , Equation 1 - 4 Equation 1 - 5 where is the free energy change per unit volume of solid formation, is the solid/liquid interfacial free energy, is the wetting angle, is the cluster free energy, is the total number of atoms, k is Boltzmans constant, and T is temperature. T herefore, the critical nucleus radius is dependent on and , while the energy barrier depends on , , and . In the case of unmodified Al - Si eutectic, the random motion of atoms within the liquid results in small unstable clusters of Si atoms being formed. The transformation from unstable clusters to stable nucle i requires two conditions. First, a decahedron bound ed by low energy facets of 111 planes is formed to minimize the surface energy resulting from the cluster formation. Second, additional undercooling supplies a driving force to overcome the free energy barrier. However, the nucleating site of the eutectic Si is still under discussion. Elliot [ 25 ] proposed that Si nucleates on nucleation centers within liquid, whereas Dahle [ 50 ] suggested that the eutectic Si could nucleate on Al dendrite tips or the nucle i within the interdendritic liquid. 2) Nucleation Kinetics of Sr - modified Si Crystals The addition of trace elements increa ses the difficulty of the nucleation of the Si crystal s . Solidification is therefore suppressed to lower temperatures. As mentioned above , the free energy barrier depends on the interfacial energy of liquid and the undercooling. It is reported 15 [ 62 ] that t he addition of Sr decreases the interfacia l energy between the eutectic liquid and solid Si , lowering the energy barrier and thus aiding nucleation of eutectic Si particles. In contrast , the addition of Sr decreases the number of Si - Si covalent bonds, reduces the number of Si nuclei in the melt, a nd increases the undercooling , hindering the nucleation of eutectic Si [ 63 ] . This effect dominates duri ng the eutectic reaction [ 62 ] . Although t his study can explain the increased undercooling temperature, modified Si morphology, and improved ductility with Sr addition, it is not clear whether the nucleation or growth process es controls the evolution of the Al - Si eutectic. In order to comprehensively understand modification mechanism of hypoeutectic Al - Si alloys with Sr addition s , both nucleation and growth processes need to be considered. 3) Growth Kinetics of Eutectic Si Crystals T he flaky morphology of Si crystals in Al - Si eutectic colonies is dependent on both thermodynamic and kinetic parameters. There are two types of heat during solidification [ 20 ] : the specific heat o f the liquid and the latent heat of fusion. The total interface undercooling ( T) of regular eutectics consists of thermal, curvature, and kinetic undercooling . Considering both the diffusion in the liquid and the thermodynamics of solid - liquid interfaces, the total interface undercooling is related to the growth velocity (V) and the interparticle spacing ( ). This relationship of the regular eutectic growth behavior was set up by Jackson and Hunt [ 64 ] for an isothermal interface and it can be shown that , where and are alloy constants determined from thermodynamic data. Three equations, calculated by differentiating the undercooling equation and setting , describe the growth mechanisms of the regular eutectic: Equation 1 - 6 Equation 1 - 7 Equation 1 - 8 16 where is the lamellar spacing at the extremum undercooling point, which is presumed to be the operative point of the regular eutectic system. However, this model cannot properly predict the anisotropic growth behavior of irregular eutectics because of the non - isotherma l interface. Magnin and Kurz [ 65 ] developed a model for irregular eutectic sys tem by setting up a non - isothermal coupling condition over the whole interface and a morphological criterion for the branching behavior of the faceted phase. The growth behavior of the irregular eutectic is shown as follows: Equation 1 - 9 Equation 1 - 10 Equation 1 - 11 where is the ra tio of me an to extremum spacing. Thus, the growth velocity of faceted Si crystals increases exponentially with undercooling. Guzik and Kopycinski [ 66 ] developed a new model that modified the MK model to consider the effect of material constants on the interlamellar spacing and protrusion of the faceted phase. In this modified model, the characteristic depression of the nonfaceted phase and the protrusion of the faceted phase can be calculated. Material constants , such as the wetting angle, diffusion coefficient, and Gibbs - Thomson coefficient , play i mportant role s in the irregular eutectic growth. All the modified models, however, are mostly focused on the two - dimensional parameterization of the irregular eutectic. The exact growth mechanism cannot be fully understood unless a mathematical model can be developed with three - dimensional parameterization. Nevertheless, the gr owth kinetics provide correlations between undercooling, growth velocity, and interparticle spacing, which are convenient for establishing the growth mechanism of eutectic Si crystals with and without trace element additions. 4) Growth Mechanism of Eutecti c Si Crystals 17 The Si flakes in Al - Si eutectic colonies promote crack initiation and propagation and hence result in low ductility. The twin plane re - entrant growth edge (TPRE) mechanism is one of the established growth mechanisms of flaky Si , which was fi rst proposed by Hamilton and Seidensticker [ 67 ] to explain the growth behavior of germanium dendrites. Kobayashi and Hogan [ 68 ] confirmed the growth of Si in Al - Si eutectic alloys by multiple twinning through the TPRE mechanism from the experimental evidence. The surface of t he solid - liquid interface on the Si 111 plane is atomically smooth, and the equilibrium Si crystal is bounded by 111 habit planes with two parallel twins. Compared with 111 flat surfaces and the re - entrant ridge with the external angle of 219 ° , the re - entr ant corner with an external angle of 141 ° is the most favorable site for Si crystals to nucleate, hence the growth of eutectic Si along directions take place rapidly at the reentrant corner. Silicon crystals with two twins can generate more re - entran t corners to continue growing. Once a triangular corner bounded entirely by ridges is formed, the Si growth stops due to the absence of any 141 ° extenal angles. Recently, Fujiwara et al. [ 23 , 69 ] modified the growth model of Si crystal and claimed that Si crystal growth continues with the two new 141 ° extenal angles formed on the growth surface of a 219 ° external angle. According to these experimental results, there are two preferential growth directions of Si crystals: and . In the case of directions, the Si can grow alternately at the re - entrant corners with a 141 ° extenal angle, while th e Si crystal along directions can grow at both the re - entrant corners with 141 ° extenal angles and the re - entrant ridges with 219 ° extenal angles. Therefore, Si tips along directions become wider during growth, while these along directions become narrow. In this case, the faceting Si crystals in eutectic colonies grow independently and form flaky structures. 5 ) Growth Mechanism of Sr - modified Eutectic Si Crystals 18 With the Sr addition, the morphology of Si crystals is modified to form a fibro us structure and thus improves ductility. According to the TPRE mechanism, in the unmodified Al - Si system there are two different restricted growth theories for Al - Si eutectic colonies modified with trace elements. One of them [ 26 , 68 ] refers to poisoning the re - entrant edges. Sr can sit at re - entrant edges, acting as a deactivator to inhibit Si growth rather than acting as a nucleation site to promote the growth process. Rather than the Si growing along directions anisotropically , isotropic growth along multiple directions takes place for the modified Al - Si eutectic colonies. Lu and Hellawell [ 70 ] claimed that Sr is rejected into the liquid ahead of the solidification front because of the extremely low solubility of Sr in both aluminum and Si . Therefore, Sr is absorbed at trace levels onto the Si layers, wh ich can induce twinning by poisoning the growing ledges of Si . Lu and Hellawell also proposed that 1.65 is the ideal ratio of trace element to Si atom radii to induce twinning. Considering both unmodified and modified Al - Si eutectics, there are still two l imitations corresponding to the experimental results that cannot be well explained by this model: the undercooling temperature of the Al - Si eutectic reaction is increased with Sr addition and eutectic structure without Sr addition is modified by a relative ly high solidification rate. Nevertheless, it is generally accepted that Sr acts as deactivator to inhibit Si growth during solidification. Therefore, the morphology of eutectic Si crystals with Sr additions is fibrous in form, which improves the ductility of Al - Si casting alloys. Deciphering the controlling mechanisms exerted by Sr over nucleation and growth is one of the central challenges , as it is crucial for understanding how fibrous Si crystals in Al - Si eutectic colonies are formed. This section has f ocused on two aspects that are central to the identification and analysis of important trends . Sr at trace levels decreases the number of Si - Si bonds that act as the nuclei of Si crystals in eutectic colonies. Hence, Sr hinders the nucleation process. Addi tionally, Sr sits at re - entrant edges of Si crystals, acting as a deactivator to inhibit 19 crystal growth. These two mechanisms provide a framework for relating the interactions between the modifier and the morphology of the eutectic colonies. Thus, the des ired mechanical properties of Al - Si based alloy can be engineered. 1.3 Development of Cast Aluminum Alloys Using Nano - structured Chemicals 1. 3 .1 Overall Goal The goal of the proposed work is to explore the role of nanostructured silanols and their effect on the solidification kinetics, microstructure, and mechanical properties of Al - Si based casting alloys . 1.3 .2 Motivation One of the major strategies to meet the new Corporate Average Fuel Economy (CAFE) standards for 2025 (54.5 mpg) [ 71 ] is to reduce vehicle weight by replacing steel with lightweight materials. The ultimate objective is to enhance aluminum alloys w ith both high ductility and strength . However, conventional Al - Si based casting all oys show low ductility as a compensation for high strength, since the presence of irregular flakes like eutectic Si crystals in the as - cast condition promotes crack initiation and propagation. To improve their overall mechanical behaviors , grain refiner s or chemical modifier s ha ve been added into the base alloy, which enhances the overall mechanical properties through microstructure refinement. For example, titanium (Ti)/ boron (B) can react with Al to form intermetallic compounds (IMCs) [ 72 ] , such as Al 3 Ti and TiB 2 . These IMCs serve as heterogeneous nucleation sites for primary Al, resulting in refined grain s ize. The addition of strontium (Sr) can poison the re - entrant corner or growing ledge s in eutectic Si in order to modify the morphology of eutectic Si from flaky to fibrous form [ 62 , 70 ] . However, as the Si content is greater than 3%, the effect of the Al - Ti - B grain refiner on changing Al grain size diminishes significantly [ 39 , 73 ] , which result s from the formation of TiSi and TiSi 2 . Sr also suffers from an accelerated oxidation at 20 casting temperature s [ 56 ] , leading to the loss of Si modification in Al - Si alloy after the second melt. Therefore, the issues of Ti poisoning and Sr fading after the addition of refiners/modifiers become major concerns for aluminum foundries and researchers. Thus, there is a need to develop new Al - based casting alloys with high ductility while not sacrificing strength for stru ctural applications without the poisoning and fading effects of current methods. This thesis relates to the use of nano - structured chemicals in the processing of lightweight aluminum alloys for structural applications. Polyhedral oligomeric silse squioxane (POSS) are silse s quioxane - based nano - structural chemicals , as shown in Figure 1 - 6 (a) [ 74 ] . These chemicals are cage - like structures with repeating monomer units of RSiO 1.5 , where Si is the element silicon, O is oxygen , and R is a hydrocarbon group, e.g. ethyl , isobutyl, phenyl , etc. These hy brid structure links inorganics to organics with unique properties, such as high reactivity and thermal stability [ 74 , 75 ] . Trisilanol POSS (TSP) are produced by replacing a Si atom at the corner of the POSS to form silanol s, as sh own in Figure 1 - 6 (b) . These silanol s have been revealed to form thermodynamically s table Si - O - M bonds with metals [ 75 ] . Silanols in TSP can react with reactive metals to make a salt and hydrogen. 2X SiOH - + M X + X(SiO) 2/X M + XH 2 Equation 1 - 1 21 Figure 1 - 6 Chemical Structure of (a) Polyhedral Oligomeric Silsesquioxane (POSS) and (b) T risilanol POSS. (a) ( b ) 22 Many metallic elements, such as Al [ 76 , 77 ] , Ti [ 76 , 78 ] , V [ 78 ] , Zn [ 77 ] , have been incorporated with TSP to form polyhedral oligomeric metallic silsesquioxane (POMS ) . Effect of TSP additions on the microstructure and mechanical reliability of tin - based solder joints has been investigated by Lee [ 79 ] and Choudhuri [ 80 ] . Their study suggested that TSP additions refine d the as - solidified microstructure of SAC305. Optical micrographs comparing the effects of 2wt% TSP additions on the bulk microstructure of SAC30 5 have been presented in Figure 1 - 6 . TSP addition result ed in sm aller sized dendritic arms as compared to the control specimen. TSP additions were also found to influence the morphology and size of IMCs (Ag 3 Sn and Cu 6 Sn 5 ) present in microstructure of SAC305. Figure 1 - 7 ( a ) shows the backscattered electron image (BEI) t ernary eutectic region in the as - solidified microstructure of SAC305, while Figure 1 - 7 ( b ) present the BEI of the ternary eutectic regions in SAC305+TSP. In SAC305, large (~30µM) rod and needle shaped Cu 6 Sn 5 and Ag 3 Sn IMCs are observed in the ternary eute ctic region, while specimens containing TSP show a refined ternary eutectic microstructure with relatively spheroidal morphology. In general addition of TSP results in a more uniform microstructure. This suggests that such addition influences the growth o f IMCs within the liquid phase during solidification. 23 Figure 1 - 6 Optical micrographs showing the effects of 2wt% TSP additions on the as - solidified microstructures of SAC305. [80] 24 Figure 1 - 7 High magnification BEI showing sizes and mo rphology of Cu6Sn5 and Ag3Sn IMCs in the ternary eutectic regions in (a) SAC305, and (b) SAC305 + 2wt% TSP .[80] (a) (b ) 25 Figure 1 - 8 Comparison of residual shear strengths of SAC305 and SAC2iB solder joints. Error bars show 95% confidence intervals. Second order exponential non - linear least square fits have also been shown. [80] The refined IMCs and SDAS with TSP addition results in t he residual mechanical improvement of SAC305 solder joints within the number of TMF cycles chosen as shown in Figure 1 - 8 . Such an improvement indicates that TSP reinforcements retard the microstructural damages responsible for the deterioration in the mec hanical strength of control joints. In summary, TSP additions refine microstructure of both the primary Sn and the Cu6Sn5 intermetallic compounds (IMC). More importantly, TSP prevents coarsening at elevated 26 temperature for microstructural stability of Sn - based solders, and also significantly improves the thermomechanical fatigue life of solder joints by pinning Sn grain boundaries from sliding and enhancing the IMC/Sn bonding . The objectives for this thesis are to achieve microstructure refinement and stability with TSP additions , leading to significant mechanical property enhancement for cast Al alloys. Presence of silanol groups in TSP facilitates the bonding between TSP and Al. Once this bonding is achieve d to form a closed Si - O - M cage, the cage structure can be maintained a temperature higher than 1000 o C without thermal degradation due to the inert nature of POSS inorganic core. Because of these unique feathers, poisoning and fading issues in cast Al alloy s with the incorporation of TSP would not occur. 1.3 .3 Fun damental Idea and Specific Aims A nanostructured chemical based on trisilanol p olyhedral silse s quioxane (TSP ) was used as a chemical modifier to modify the microstructure of Al - Si based casting alloys. This study focuses on increasing the ductility of Al - Si casting alloys by modifying the microstructure of eutectic Si crystals in the as - cast condition while maintaining the strength of the unmod ified alloy. Specific aims of this project include the following: 1) Study the impact of TSP additions on the microstructure and mechanical properties of Sn based binary and ternary alloys. The relationships between the structure and properties of the mate rial were investigated to design microstructure with a predetermined set of material properties. 2) Investigate the effect of TSP additions on the solidification of Al - Si casting alloys to understand how their microstructures are produced during phase tran sformation. The role of phase transformations in the microstructure development of TSP modified Al - Si w ere addressed to propose the modification mechanism . 27 3) Develop an optimized process to obtain a TSP master modifier to avoid the fading issues of curren t modifiers. By controlling the TSP incorporation method and TSP content in the master , the designed master modifier is demonstrated to function in hypoeutectic Al - Si for structural applications . 4) Use the TSP master alloy to m odify the microstructure of Al - Si binary and commercial alloys with improved mechanical properties . Microstructural and thermal characterization s have been performed to ascertain the role of the TSP master alloy in the strength and ductility changes. This study addresses fundamental questions on how the addition of TSP influences microstructure and solidification behavior of Al - Si based casting alloys and provides solutions to reinforce aluminum alloys with high ductility and strength for automotive industry. 28 Chapter 2 Sn - Cu Binary Alloys Refinement 2.1 Background Tin (Sn) based lead - free alloys have been widely used as industrial solders to fabricate solder joints in recent years. [ 81 ] Tin - silver - copper (SAC) alloys are considered as the preferred substitute fo r Sn - Pb solder in solder joint applications. [ 82 , 83 ] However, Sn - Cu binary alloys, such as Sn - 0.7wt.%Cu and Sn - 3wt.%C u, have been drawn more and more attentions as the prevailing alloy system due to slightly higher service temperature and lower cost as compared to SAC solder alloys. The microstructure of Sn - Cu alloy is determined by the chemical composition of Cu in Sn. In the case of eutectic Sn - 0.7Cu, the microstructure is typically formed with intermetallic compound (IMC) particles embedded in Sn - rich matrix, while hypereutectic Sn - Cu contains proeutectic Cu 6 Sn 5 phase besides eutectic constituent. Unfortunately, because of the short thermo - mechanical fatigue lifetime, Sn - Cu alloys have poor strength and creep resistance [ 84 ] , and hence limits their solder joint applications. The limiting mechanical performances are mainly due to the size and texture of Sn - rich grains, the size and dispersion of Cu 6 Sn 5 IMC during processing [ 85 , 86 ] , and coarsening of IMC during service [ 87 , 88 ] . Currently, there are two major approaches to solve these problems: alloying and composite approaches. For alloying approach, it has been reported that N i [ 89 - 91 ] , Co [ 89 ] , A l [ 92 ] , Zn [ 93 , 94 ] , In [ 95 ] , and/or rare earth element s [ 96 ] are added as grain refiner of Sn - rich grains. The potential issue is that the addition of minor alloying elements cannot prevent IMC coarsening during services. For composite approach, additional IMC are incorporated by either in - situ or mechanical mixing methods . [ 97 , 98 ] The drawbacks of IMC incorporation are the big size of reinforcements (several microns) resulting in poor mechanical p roperties, as well as coarsening of reinforcements during service. Another composite approach is to incorporate inert reinforcements by mechanical mixing method. [ 99 , 100 ] Lack of bonding between solder and 29 reinforcements [ 101 ] , clustering of particles with pores in - between agglomerated particles, and interfacial cracking due to the required mechanical work to break up clusters are the unavoidable issues. Polyhedral oligomeric silsequio xanes (POSS) are silsequioxane - based nano - structural chemicals. [ 74 , 75 ] These chemicals are cage like structures with repeated monomer units of RSiO 1.5 where Si is the element silicon, O is oxygen, and R is hydrocarbon group, e.g. ethyl, isobutyl, phenyl etc. Partial cage - like POSS with multiple silanol (Si - OH) function alities are produced by replacing a Si atom at the corner of POSS to form silanol bonds. TriSilianol phenyl POSS (TSP) is used in this study. These silanol bonds have revealed to form thermodynamically stable Si - O - M bonds with metals (e.g. Sn). The object ive of this study is to investigate creep resistance and strength, and correlate above mentioned mechanical properties changes with microstructure alterations of Pb - free Sn - Cu solder alloys with addition of nanostructured silanols. 2.2 Experimental 2.2.1 P reparation of Sn/Cu Solder Wires TriSilanolphenyl Polyhedral Oligomeric Silsesquioxane (TSP), obtained from Hybrid Plastics, Inc. (Hattiesburg, MS), was used in this study as nano - structured modifier. Nearly eutectic, Sn - 0.7Cu, and hypereutectic, Sn - 3Cu, alloys were manufactured by Johnson Manufacturing with their standard industrial alloy practice. For a comparative study of creep resistance and microstructures of TSP - treated Sn - Cu alloys, both Sn - 0.7Cu and Sn - 3Cu alloys containing 0.5 wt% TSP were prepared using the same alloying approach. Nearly eutectic and hypereutectic Sn - Cu alloys wit hout and with 1wt% TSP addition were cast to 38 mm by 178 mm billet form according to the established ISO procedures by Johnson 30 Manufacturing. 0.8 mm diameter solid solder wires were then extruded from billet using a 100 Ton Collin Press. 2.2.2 Coil Creep Testing The creep behavior of solder wires was characterized by coil creep test. A 20 - turn coil of Sn - Cu solder wire was wound on a cylindrical aluminum tube with 25.4 mm in diameter as shown in Figure 2 - 1 ( a ) . The creep test of a self - loaded coil in Figu re 2 - 1 ( b ) was conducted in a temperature - controlled column at 40, 50, 60, and 70°C. Stretch rate was evaluated through dividing the spacing between the 1 st and the 10 th turns by time at specific temperature for all wires, which can be used as an indicator of creep rate. Figure 2 - 1 Creep E xperimental S etup: (a) S tereo M icrographs of W ire W inding and (b) S chematic of C oil C reep S etup . 2.2.3 Polarized Optical Microscopy Four pieces of each specimen were cut from four different regions and mounted with epoxy resins in such a way that wires orientated in the extrusion direction. The wire mount was ground 31 until wire surface reached the center, and then polished. Polarized op tical micrographs were taken using an inverted metallographic microscope, Nikon ECLIPSE MA200. 2.2.4 Differential Scanning Calorimetry Thermal characteristics were performed using a TA Instrument Q2000 equipped with a mechanical cooling system under N 2 at mosphere. Sn - Cu wires, about 15mg, were placed in a TZero TM aluminum pan made for thermal analysis of solid samples. All samples were heated to 350°C and kept for 1 min to ensure complete melting, then cooled to 150°C at a constant cooling rate of 10°C/min and then reheated to 350°C at a heating rate of 10°C/min. Onsite of the exothermic solidification peak was denoted as the crystallization temperature, T c , and the onsite of the endothermic melting peak was denoted as the melting temperature, T m . 2.2.5 Ind entation Microhardness Microhardness of the polished wire in the epoxy mount was measured with a Cl ark CM - 800 Knoop tester at 25 gf applied load with 15 seconds dwell at room temperature. 2.2.6 Orientation Imaging Microscopy The Sn crystal orientation of the polished wire in the epoxy mount was evaluated using a CamScan field - emission gun (FEG) secondary electron microscopy (SEM) equipped with electron backscatter diffraction (EBSD) analysis systems. The accelerating voltage used was 20 kV. A 3 - e was used for EBSD scans. To reduce the charging effects inside the SEM from the nonconductive epoxy mount, the mounted sample was mostly covered by Cu tape, and the areas surrounding the wires were covered with conductive carbon paint. EBSD c - axis orient ation maps, texture, and pole figure plots were generated using orientation imaging microscopy (OIM) Data analysis software. 32 2.3 Results and Discussions 2.3.1 Creep Behavior The addition of TSP improved creep resistance of both Sn - 0.7Cu and Sn - 3Cu wires. Stretch rates of both Sn - 0.7Cu and Sn - 3Cu solder wires without and with TSP addition at various temperatures were plotted in Figure 2 - 2 . It is well known that there are typically three stages with distinctive strain - time feather during creep: primary, ste ady - state, and tertiary creep. In the solder coil study, steady - state creep dominates the creep behavior due to the time - independent strain rate at all testing temperatures. Diffusion creep and dislocation creep are basically two general mechanisms by whic h steady - state creep occurs. The melting points of eutectic Sn - Cu alloy was around 500 K [ 102 ] , so the homologous temperature at room temperature (300K) is approximately 0.6. The homologous temperatures during coil creep testing were greater than 0.6. Therefore, the dominate creep mechanism is diffusion creep. Comparing Figure 2 - 2 ( a ) with Figure 2 - 2 ( b ) , without TSP addition, Sn - 3Cu solder wire had a higher creep resistance than Sn - 0.7Cu. The higher creep resistance of Sn - 3Cu over Sn - 0.7Cu suggested that the diffusion creep behavior in Sn - Cu alloys was controlled by the microstructure and spatial distribution of Cu 6 Sn 5 IMCs, similar to precipitation hardening [ 103 ] . The results in this study were consistent with the work done by Subramanian et al. [ 101 ] , where in - situ Cu 6 Sn 5 IMCs reinforced eutectic Sn - Ag showed better creep resistance than the control eutectic Sn - Ag at room temperature because of the longer t ime for a dislocation to climb during diffusion creep. In Figure 2 - 2 , the stretch rate of both Sn - Cu wires at four different temperatures was reduced due to the addition of TSP. More importantly, effect was more significant for Sn - 0.7Cu. Thus, Sn - Cu alloys with TSP addition exhibited higher resistance to creep, especially for eutectic Sn - 0.7Cu solder wire. This improved creep resistance can be explained by the microstructure change of Cu 6 Sn 5 IMCs with TSP addition. 33 Figure 2 - 2 Stretch Rates at Various Temperatures: (a) Sn - 0.7Cu Wire without and with TSP Additions and (b) Sn - 3Cu Wire without and with TSP Additions. 34 2.3.2 Microstructure TSP addition to the solder refined the intermetallic compound in the solder matrix. The microstructure of eutectic Sn - 0.7Cu solder was shown in Figure 2 - 3 ( a ) with dot - shaped Cu 6 Sn 5 IMC particles embedded in Sn - rich matrix, whereas the microstructure of hy pereutectic Sn - 3Cu solder was shown in Figure 2 - 3 ( b ) with proeutectic rod - shaped Cu 6 Sn 5 phase and eutectic constituent. Figure 2 - 2 showed that the creep resistance of Sn - 3Cu wire was much higher than Sn - 0.7Cu. This can be explained by the presence of the large and rod - shaped IMCs in Figure 2 - 3 ( b ) , which hindered dislocation climb and improved the steady - state creep. Under the same wire extruding conditions, Cu 6 Sn 5 IMC particles in both TSP - treated Sn - Cu alloys appeared to be smaller and well distributed as compared to control, as shown in Figure 2 - 3 (c) and 2 - 3 ( d ) . TSP - treated Sn - Cu alloys with smaller IMCs and closer spacing between IMCs exhibited higher resistance to creep. Cu 6 Sn 5 IMC particles size in Sn - 3Cu was larger than those in Sn - 0.7Cu, resulti ng in less improvement of creep resistance for Sn - 3Cu with TSP incorporation as compared to Sn - 0.7Cu. The addition of TSP also caused a refined Sn - rich morphology in eutectic Sn - 0.7Cu. Comparing Figure 4a with Figure 4c, the Sn - rich grain size in TSP - trea ted Sn - 0.7Cu was smaller than that in Sn - 0.7Cu control. Based on how paths change during diffusion creep, diffusion creep can be classified into two main classes: Nabarro - Herring creep and Coble creep, where the diffusion paths in Nabarro - Herring creep are predominantly through the main body of the grains at relative high temperature, and the diffusion paths in Coble creep are predominantly through the grain boundaries at relative low temperature. With TSP addition, both the finely dispersed IMC particles w ithin Sn - rich grains and the refined Sn - rich grains improved creep resistance, especially for Sn - 0.7Cu, suggesting that the dominance of diffusion creep in Sn - Cu system is supported by Nabarro - Herrring creep. Nevertheless, the improved creep resistance for both Sn - 0.7Cu and Sn - 35 3Cu was attributed to the smaller IMCs and closer spacing between IMCs with TSP incorporation. This changed microstructure can be explained by thermal analysis. 2.3.3 Solidification of Eutectic Constituent The addition of TSP reduce d the extent of undercooling for eutectic solidification reaction in both Sn - 0.7Cu and Sn - 3Cu. Effect of TSP addition on the eutectic solidification reaction of Sn - 0.7Cu and Sn - 3Cu were examined using Differential Scanning Calorimetry (DSC). Table 2 - 1 lis ted the values of the crystallization temperature, Tc, the melting temperature, Tm, and the m denoted erence between Tm and Tc. The melting temperature of eutectic constituent in Sn - 0.7Cu and Sn - 0.3Cu appeared at 226°C. The melting temperature of eutectic constituent in both alloys did not change with TSP addition. The extent of undercooling, which was det ermined by the difference between the melting temperature and the cooling temperature, was used to quantify the solidification behavior. Table 2 - 1 Comparison of S olidification T emperature, M elting T emperature , and U ndercooling for Sn - Cu W ires . T c ( °C ) T s ( °C ) Tm ( °C ) °C ) Sn - 0.7Cu - 202.1 226.5 24.3 TSP - treated Sn - 0.7Cu - 20.8 226.5 21.7 Sn - 3Cu 318.9 198.4 226.3 27.9 TSP - treated Sn - 3Cu 313.1 217.3 226.3 8.9 *Tc: formation temperature of the proeutectic constituent Ts: solidification temperature of the eutectic constituent Tm: melting temperature of the eutectic constituent for eutectic solidification reaction , Tm - Tc Based on classical nucleation theory, the additional undercooling supplies a driving force to overcome free energy barrier during solidification reaction. In Table 2 - 1 , the undercooling of 36 Sn - 0.7Cu and Sn - 3Cu was around 24 and 28, respectively. These resul ts suggested that eutectic solidification reaction was not affected by the presence of proeutectic Cu 6 Sn 5 IMCs, as Sn - 0.7Cu is a near eutectic, while Sn - 3Cu is hypereutectic. However, the undercooling of TSP - treated Sn - Cu solder alloys for eutectic reactio n was reduced in both. These results suggested that TSP addition decreased the interfacial energy between solid and liquid, lowered the energy barrier, and thus aided nucleation of IMCs as nucleation sites during eutectic reaction. Effect of TSP addition o n eutectic reaction in Sn - 3Cu was significant, which was consistent with results of optical microscopy showing presence of micron - size Cu 6 Sn 5 in Sn - rich grains. That indicated that TSP could act as heterogeneous nucleation sites to improve nucleation rate of eutectic reaction and affect the size and spatial distribution Cu 6 Sn 5 IMCs. Figure 2 - 3 Polarized O ptical M icrographs of Sn - Cu S older W ires: (a) Sn - 0.7Cu W ire, ( b ) TSP - treated Sn - 0.7Cu W ire, ( c ) Sn - 3 Cu W ire, and (d) TSP - treated Sn - 3Cu W ire . 37 2.3.4 Indentation Microhardness The addition of TSP improved microhardness of both Sn - 0.7Cu and Sn - 3Cu wires. Knoop hardness values of solder wires made of both Sn - 0.7Cu and Sn - 3Cu without and with TSP addition we re shown in Figure 2 - 4 . Similar hardness values were observed between Sn - 0.7Cu and Sn - 3Cu, suggesting hardness of Sn - Cu is more sensitive to the microstructure of Sn - rich grains than IMC distribution. Hardness of Sn - 3Cu wire increased from 11.9 to 13.1 H k with TSP incorporation (about 12% improvement). Increase in hardness due to TSP addition in Sn - 0.7Cu was about one - half of that in Sn - 3Cu. It is clear from indentation microhardness test that the addition of TSP enhances the microhardness of both Sn - 0.7Cu and Sn - 3Cu wires. These results were also consistent with our previous study on eutectic Sn - Ag [ 75 ] . The observed improvement in strength of TSP - treated Sn - Cu wires was further explained by the microstructure change of Sn - rich grains. Figure 2 - 4 Knoop H ardness V alues of Sn - Cu W ires . 38 2.3.5 EBSD Results on Tin - rich Grains TSP addition showed significant effect on the refinement of Sn - rich phase. Figures 8a and 8b were the EBSD Sn crystal c - axis orientation map of solder wires made of Sn - 0.7Cu without and with TSP addition. Sn - 0.7Cu wire, as shown in Figure 2 - 5 ( a ) , exhibited a non - uniform grain size distribution due to extrusion process. Under the same extruding process, S n - 0.7Cu wire with TSP addition in Figure 2 - 5 ( b ) exhibited a more uniform grain size distribution as compared to control. The EBSD Sn c - axis orientation maps of solder wires made of Sn - 3Cu without and with TSP addition were shown in Figure 2 - 5 ( c ) and 2 - 5 ( d ) , respectively. Sn - 3Cu wire exhibited a more uniform grain size distribution as compared to Sn - 0.7 due to the presentence of proeutectic Cu 6 Sn 5 phase. However, average size of Sn - rich grains in Sn - 3Cu was reduced with TSP incorporation. 39 Figure 2 - 5 C - axis orientation maps of Sn - Cu solder wires: (a) Sn - 0.7Cu wire, (b) TSP - treated Sn - 3Cu wire, (c) Sn - 3Cu wire, and (d) TSP - treated Sn - 3Cu wire These results corresponded well with the Hall - Petch equation which gives the relationship y and grain size: y 0 + k y d - 1/2 (1) 40 0 and k y are constants. Therefore, the reduced grain size can be explained by the improvement of strength with TSP incorporation. The reinforcement mechanism of lead - free solder by TSP in previous stud y [ 75 ] was proposed that t he improvement of thermomechanical fatigue performance results from the grain boundary pinning effect of TSP. In this study, effect of TSP on the strength of Sn - Cu solid solder wire also supports the proposed reinforcement mechanism. 2.3. 6 Texture of Tin - rich Grains The texture intensity of both Sn - 0.7Cu and Sn - 3Cu wires was reduced with TSP incorporation. Both discrete and density pole figures in Figure 2 - 6 were plotted from the Sn c - axis orientation maps in Figure 2 - 5 . The pole figures in Figure 2 - 6 ( a ) showed that observed orientations were represented by a cluster of poles near upper left of the (001) discrete pole figure and a maximum value of 7.079 in the (001) density pole figure, indicating a preferred crystal orientation. With TSP addition, the te xture intensity was reduced from 7.079 to 6.042, as shown in Figure 2 - 6 ( b ) . Similarly, the pole figures in Figure 2 - 6 ( c ) showed that observed orientations were represented by a cluster of poles near lower left of the (001) discrete pole figure and a maxi mum value of 10.343 in the (001) density pole figure, indicating a strongly preferred crystal orientation. With TSP addition, the texture intensity was reduced from 10.343 to 6.314, as shown in Figure 2 - 6 ( d ) . The reduced texture intensity of both Sn - 0.7Cu and Sn - 3Cu wires can be attributed to the changed microstructure, size, and spatial distribution of Sn - rich grains and Cu 6 Sn 5 IMCs with TSP incorporation. 41 Figure 2 - 6 (001) and (100) discrete and density pole Fig.s of Sn - Cu solder wires: (a) Sn - 0.7Cu wire, (b) TSP - treated Sn - 3Cu wire, (c) Sn - 3Cu wire, and (d) TSP - treated Sn - 3Cu wire 2.4 Conclusions In summary, sets of experiments were carried out on both nearly eutecti c Sn - 0.7wt%Cu and hypereutectic Sn - 3wt%Cu solid solder wires to identify the effect of nanostructured silanols addition on creep resistance and microstucture of Sn - Cu solder wires. The following conclusions can be made from the current work: 42 1. The extent of undercooling for eutectic solidification reaction was reduced for both Sn - Cu wires with TSP addition, affecting the microstructure, size and spatial distribution, of Sn - rich grains and Cu 6 Sn 5 IMCs. 2. Addition of TSP improved uniformity of Sn - rich grains and reduced the texture intensity of Sn - rich grains of both Sn - Cu wires. 3. Addition of TSP reduced size and improved spatial distribution of Cu 6 Sn 5 IMCs in both Sn - Cu wires. 4. The decreasing size of Sn - rich grains was due to the grain boundary pinning effect of TSP, whereas nano - sized cage structure of - Si - O - Sn served as heterogeneous nucleation sites for Cu 6 Sn 5 IMCs to reduce IMC size and improve IMC distribution. 5. These microstructure changes in Sn - 0.7Cu significantly improved its creep resistance, with minima l increase in hardness, while in Sn - 3Cu significantly increased the hardness, with minimal increase in creep resistance. 43 Chapter 3 Sn - based Ternary Alloys 3.1 Background To further improve mechanical properties of cast alloys , the evolution of secondary phases, su ch as precipitates or IMCs, also needs to be fully optimized. Because some properties, such as strength, stiffness, creep, and fatigue, strongly depend on the microstructure of secondary phase. In previous study, TSP is added into Sn - Cu binary alloys to im prove creep re sistance of solid solder wires. The improved creep behavior is due to the controllable size and distribution of IMCs in Sn - Cu alloys with TSP additions. Sn - rich ternary alloys have attracted considerable attention due to their good solderability and wetting properties. It is important to inverstiage how TSP additions affect their as - cast microstructure as well as creep resistatance. Many Sn based ternary alloy systems with different alloying elemnts such as Ag, Cu, Zn, and Sb have been developed. The secondary creep state or steady state is of the greatest engineering importance since in this regime plastic flow will occur while the stress and temperature are held constant and the strain rate (d /dt) dependence can be expressed as the following equation describes: equation 3 - 1 w valu e of stress exponent, n, is related to operative creep mechanisms, while the activation energy, Q, correlates to the activation energy for diffusion. 44 3.2 Experimental 3.2.1 P reparation of Solid Wires TriSilanolphenyl Polyhedral Oligomeric Silsesquioxane ( TSP), obtained from Hybrid Plastics, Inc. (Hattiesburg, MS), was used in this study as nano - structured modifier. Pure tin and three ternary alloys were manufactured by Johnson Manufacturing with their standard industrial alloy practice as shown in Table 3 - 1 . For a comparative study of creep resistance and microstructures of 0.5% TSP - treated alloys were prepared using the same alloying approach. All alloys with and without 0.5 wt% TSP addition were cast to 38 mm by 178 mm billet form according to the establish ed ISO procedures by Johnson Manufacturing. 0.8 mm diameter solid solder wires were then extruded from billet using a 100 Ton Collin Press. Table 3 - 1 Chemical Composition in wt% of All Wires. Material Cu Ag Sb Zn Sn Pure Sn - - - - 99.9 Sn - 4.7Ag - 1.7Cu 1.7 4.7 - - Balance Sn - 3.5Cu - 7.5Sb 3.5 - 7.5 - Balance Sn - 15Cu - 3Zn 15 - - 3 Balance 3.2.2 Coil Creep Testing The creep behavior of solder wires was characterized by coil creep test. A 20 - turn coil of solder wire was wound on a cylindrical aluminum tube with 25.4 mm in diameter as shown in Figure 2 - 1 ( a ) . The creep test of a self - loaded coil in Figure 2 - 1 ( b ) was conducted in a temperature - controlled column . Different temperatures were used to keep the testing temperature within 0.4 - 0.7 range of their melting temperatures . Stretch rate was evaluated through dividing the spacing between the 1 st and the 10 th turns by time at specific temperature for all wires, which 45 can be used as an indicator of creep rat e. To evaluate the stress exponent ´n´ and the activation energy ´Q´ of different alloys, points on the graph, with a linear regression, the gradient of the straight line would be n - th e stress exponent. Similarly, by plotting ln (strain rate) against the reciprocal of temperature, the gradient of the plot should be - Q/R. Using this method, activation energy of each material could be obtained from the plot. 3.2.3 Polarized Optical Micr oscopy Four pieces of each specimen were cut from four different regions and mounted with epoxy resins in such a way that wires orientated in the extrusion direction. The wire mount was ground until wire surface reached the center, and then polished. Polar ized optical micrographs were taken using an inverted metallographic microscope, Nikon ECLIPSE MA200. 3.3 Results and Discussions 3.3.1 Pure Tin The addition of TSP improved creep resistance of pure Sn wires dramatically . Stretch rates of Sn wires with an d with out TSP addition s at two temperatures were plotted in Figure 3 - 1. T he stretch rate of tin wires at two different temperatures was reduced due to the addition of TSP. As it can be observed in Table 3 - 2, the calculated stress exponents are between 2 an d 6, which is consistent with the fact that this is dislocation creep. However, TSP additions increase creep activation energy from 2.4 to 85.2 KJ/mol, resulting in high creep resistance. 46 Figure 3 - 1 Stretch Rates of Pure Sn and TSP - treated Sn Wires at Various Temperatures. Table 3 - 2 Stress Exponent and Activation Energy of Pure Sn and TSP - treated Sn Wires. Material Temperature ( o C) Stress Exponent Activation Energy (KJ/mol) Pure Sn 26 6.1 2.4 30 5.6 TSP - treated Sn 26 2.8 85.2 30 2.7 47 Figure 3 - 2 Optical Micrographs of Pure Tin Solder Wires: (a) Pure Sn Wire and (b) TSP - treated Sn Wire. Figure 3 - 2 shows the microstructure of pure Sn wires with and without TSP additions. As it is a pure material, pure Sn, all the grains that appear on the image are Sn grains with different grain size. An alyzing the modified sample, it can be observed almost no variation on the microstructure of Sn. 3.3.2 Sn - Ag - Cu Alloys The addition of TSP also improved creep resistance of Sn - Ag - Cu ternary wires dramatically . Stretch rates of Sn - 4.7Ag - 1.7Cu wires with and with out TSP addition s at four temperatures were plotted in Figure 3 - 3. T he stretch rate of tin wires at four different temperatures was reduced due to the addition of TSP. As it can be observed in Table 3 - 3, the calculated stress exponent is about 1 in all the samples, which is consistent with the fact that it is diffusion creep, being the main mechanism is thes e solder specimens over this temperature range. However, TSP additions increase creep activation energy from 62.4 to 150.7 KJ/mol, resulting in high creep resistance. (a) (b ) 48 Figure 3 - 3 Stretch Rates of Sn - 4.7Ag - 1.7Cu and TSP - treated Sn - 4.7Ag - 1.7Cu Wires at Various Temperatures. 49 Table 3 - 3 Stress Exponent and Activation Energy of Sn - 4.7Ag - 1.7Cu and TSP - treated Sn - 4.7Ag - 1.7Cu Wires. Material Temperature ( o C) Stress Exponent Activation Energy (KJ/mol) Sn - 4.7Ag - 1.7Cu 40 0.6 62.4 50 0.7 60 0.6 70 0.8 TSP - treated Sn - 4.7Ag - 1.7Cu 40 1.0 150.7 50 1.9 60 1.8 70 1.0 50 Figure 3 - 4 Optical Micrographs of Sn - 4.7Ag - 1.7Cu Solder Wires: (a) Sn - 4.7Ag - 1.7Cu Wire and (b) TSP - treated Sn - 4.7Ag - 1.7Cu Wire. Figures 3 - 4 illustrates the microstructure of Sn - 4.7Ag - 1.7Cu wires with and without TSP additions. Three ph ases were observed in Sn - 4.7Ag - 1.7Cu wire. The primary Sn phases, Cu6Sn5 and Ag3Sn intermetallic compounds. With TSP additions, both grain size of primary Sn and IMCs size decreased. This refinement with TSP additions is associated with the increased activ ation energy and creep resistance. 3.3.3 Sn - Cu - Sb Alloys The addition of TSP improved creep resistance of Sn - Cu - Sb ternary wires dramatically . Stretch rates of Sn - 3.5Cu - 7 .5Sb wires with and with out TSP addition s at four temperatures were plotted in Figure 3 - 5. T he stretch rate of Sn - Cu - Sb wires at four different temperatures was reduced due to the addition of TSP. As it can be observed in Table 3 - 4, the calculated stress (a) (b ) 51 exponent is about 1 in all the samples, which is consistent with the fact that it is di ffusion creep, being the main mechanism is these solder specimens over this temperature range. However, TSP additions increase creep activation energy from 111.6 to 123.7 KJ/mol, resulting in high creep resistance. Figure 3 - 5 Stretch Rates of Sn - 3.5Cu - 7.5Sb and TSP - treated Sn - 3.5Cu - 7.5Sb Wires at Various Temperatures 52 Table 3 - 4 Stress Exponent and Activation Energy of Sn - 3.5Cu - 7.5Sb and TSP - treated Sn - 3.5Cu - 7 .5Sb Wires. Material Temperature ( o C) Stress Exponent Activation Energy (KJ/mol) Sn - 3.5Cu - 7.5Sb 26 1.2 111.6 35 1.0 40 1.0 45 1.6 TSP - treated Sn - 3.5Cu - 7.5Sb 26 0.9 123.7 35 0.7 40 1.2 45 1.6 53 Figure 3 - 6 Optical Micrographs of Sn - 3.5Cu - 7.5Sb Solder Wires: (a) Sn - 3.5Cu - 7.5Sb Wire and (b) TSP - treated Sn - 3.5Cu - 7.5Sb Wire. Figures 3 - 6 illustrates the microstructure of Sn - 3.5Cu - 7.5Sb wires with and with out TSP additions. Three phases were observed in Sn - 3.5Cu - 7.5Sb wire. The primary Sn phases, Cu3Sn and SbSn3 intermetallic compounds. With TSP additions, both grain size of primary Sn and IMCs size decreased. This refinement with TSP additions is associate d with the increased activation energy and creep resistance. 3.3.4 Sn - Cu - Zn Alloys The addition of TSP reduced creep resistance of Sn - Cu - Zn ternary wire s . Stretch rates of Sn - 15Cu - 3Zn wires with and with out TSP addition s at three temperatures were plotted in Figure 3 - 7. T he stretch rate of Sn - Cu - Zn wires at four different temperatures was increased due to the addition of TSP. As it can be observed in Table 3 - 5, the calculated stress exponent is about 1 in (a) (b ) 54 all the samples, which is consistent with the fact that it is diffusion creep, being the main mechanism is these solder specimens over this temperature range. However, TSP ad ditions decrased creep activation energy from 118.7 to 88.7 KJ/mol, resulting in low creep resistance. Figure 3 - 7 Stretch Rates of Sn - 15Cu - 3Zn and TSP - treated Sn - 15Cu - 3Zn Wires at Various Temperatures. 55 Table 3 - 5 Stress Exponent and Activation Energy of Sn - 15Cu - 3Zn and TSP - treated Sn - 15Cu - 3Zn Wires. Material Temperature ( o C) Stress Exponent Activation Energy (KJ/mol) Sn - 15Cu - 3Zn 65 0.8 118.7 70 0.8 80 0.8 TSP - treated Sn - 15Cu - 3Zn 65 0.3 88.8 70 0.6 80 0.8 56 Figure 3 - 8 Optical Micrographs of Sn - 15Cu - 3Zn Solder Wires: (a) Sn - 15Cu - 3Zn Wire and (b) TSP - treated Sn - 15Cu - 3Zn Wire. Figures 3 - 8 illustrates the microstructure of Sn - 15Cu - 3Zn wires with and without TSP additions. Three phases were observed in Sn - 15Cu - 3Zn wire: primary Sn and (Cu,Zn)3Sn phases. With TSP additions, (Cu,Zn)3Sn size become smaller. This refinement with TSP additions is associated with the reduced activation energy and creep resistance. 3.4 Conclusions TSP additions improved creep activation energy of pure tin, Sn - Ag - Cu and Sn - Cu - Sb alloys without altering the creep deformation mechanism. The improved creep resistance with TSP additions for Sn - Ag - Cu and Sn - Cu - Sb ternary alloys is due to refined secondary phases. TSP additions also caused the refinement of the secondary phase of Sn - Cu - Zn alloys, however, the refined microstructure causes the reduction of creep resistance. (a) (b ) 57 Chapter 4 Al - 12Si Alloys with TSP Treatment 4 .1 Background Aluminum s ilicon based casting alloys have structural applications that are directly linked their light - weight and low life - cycle costs, in industries such as automotive, military, aerospace, and alternative energy [ 1 , 4 ] . However, conventional Al - Si based casting alloys exhibit low ductility as a concession for high strength, since the presence of irregular flakes like eutectic Si crystals in the as - cast condition promote crack initiation and propagation. In order to improve the ov erall mechanical performance of cast aluminum alloys , chemical modifiers , such as Na [ 45 ] , Sr [ 50 , 52 ] , Ba [ 46 ] , P [ 104 ] , and Ti/ B [ 105 ] , are added into alloys in trace levels to modify the microstructure morphology. For example, the presence of Sr changes the morphology of eutectic Si crystals from flaky to fibrous form. However, the fading and poisoning effects of Sr addition s remain two major concerns in indu st rial practice [ 56 ] . Specifically, the microstructure of a Sr - modified Al - Si alloy s cannot be maintained after the second melt, which is refe r r ed to as fading . Other trace elements have been used to improve the performance of cast aluminum alloys, i.e. Ti and B were used to refine the primary Al phase in hypoeutectic Al - Si [ 105 ] , while P was used to refine the primary Si phase in hypereutectic Al - Si [ 104 ] . However, the effectiveness of eutectic modifier s can be diminished if a hypo/ hypereutectic modifier is present, i.e. P can coarsen the eutectic structure by reacting with Sr [ 56 ] . This poisoning issue of extraneous reactions , due to the presence of trace elements in Al - Si based alloys , limits their potential in applications. Therefore, there is a need for a universal modifier that can be used in a larger range of Si content s to change the morphology of Al - Si based casting alloys and enhance their ductility. The objective in this chapter is to investigate if TSP can serve a purpose similar to the commonl 58 test this hypothesis, the TSP - treated Al - 12Si powders were melted and cast into ingots. Solidification, microstructure, and mechanical properties were evaluated. 4 .2 Experimental 4 .2.1 Materials SO1458 - t ris ilan olphenyl polyhedral oligomeric s ilsesquioxane ( Ph - TSP), obtained from Hybrid Plastics, Inc. (Hattiesburg , MS), was used in this study . The 50 µm A4047 powder , with a nearly eutectic composition (nominal wt. %) of Al - 12Si , was manufactured by Johnso n Manufacturing (Princeton, IA). 4 .2.2 Surface T reatment of A4047 powders To treat the A4047 powders with TSP, t he oxides from the powder surface s were first removed by soaking the powder in 8 wt% HCl for 5 minutes at room temperature. T he clean powers then were soaked in 1wt%TSP ethanol solutions for 24 hours to get the TSP - treated powders. 4 .2. 3 Preparation of A4047 I ngots Dried TSP - treated A4047 powders were melted at 1088 K ( 815°C ) . The molten melt was poured by gravity in to a permanent graphite mold to form a rectangular ingot of 25 mm x 50 mm x 12 mm in size and cooled in air. An A4047 ingot, which was made using the same casting approach from the clean A4047 powder without TSP treatment, was benchmarked as a control sample. To study the TSP effectiveness in A4047, the TSP - treated A4047 ingot was designated as the master alloy and was added in to untreated, molten A4047 at 1088 K (815 °C ) to get the diluted ingot with a different TSP concentration. Table 4 - 1 shows the overall composition of TSP in the A4047 ingots before and after sequential dilution. 59 Table 4 - 1 Overall C omposition of TSP in the A4047 I ngots . Samples PPM of TSP A4047 control 0 TSP - treated A4047 10,000 50% untreated A4047 + 50% TSP - treated A4047 5000 75% untreated A4047 + 25% TSP - treated A4047 2500 94% untreated A4047 + 6% TSP - treated A4047 600 98.5% untreated A4047 + 1.5 % TSP - treated A4047 150 60 4 .2. 4 Microscopy The cast ingot was cut , exposing the middle interior region of the sample, and mounted with epoxy resins . Metallographic samples were prepared by grinding with SiC paper and then 2 O ptical micrographs were taken using a Nikon ECLIPSE MA200 inverted metallographic microscope . Scanning e lectron microscopy (SEM) was performed with a LEO 1450VP SEM. 4 .2.5 Tensile Test Cylindrical tensile specimens of the Al - based alloys were machined from ingots with a gauge diameter of 4 mm and gauge length of 30 mm, according to ASTM - E8 standard for room temperature testing. 4 .2. 6 Indentation M icrohardness Microha rdness of the eutectic microconstituent of the polished sample in the epoxy mount was measured with a Cl ark CM - 800 Vickers tester using 100 gf applied load with 15 sec onds dwell at room temperature. Ten measurements for each sample were performed to obtain average and standard deviation values. 4 .2. 7 Thermal A nalysis Twelve - gram ingots of the A4047 samples were melted in a furnace at 1088 K ( 815 °C ) for 5 minutes. After the samples were i n the liquid state, the molten metal was poured into a ceramic crucible to solidify at room temperature. The ceramic crucible was connected to a Protek TM 506 digital thermometer using an N - type thermocouple wire. The time - temperature data was recorded usin g Protek TM 506 software until the alloy completely solidified. 61 4 .3 Results and Discussion 4 .3.1 Mechanical P roperties A4047 powders with and without TSP treatment were used to make ingots for mechanical property measurement. The set of the ingots was made from the same powders without TSP treatment and was benchmarked as control samples. Tensile tests were conducted on the tensile bars machined from the ingots as described earlier. The ultimate tensile strength (UTS), and elongation to fa ilure are summarized in Figure 4 - 1 . Elongation to failure increased from 5% with no TSP treatment to 18 % when treated with TSP. However, the ultimate tensile strength did not decrease with the increase of elongation. Instead, it was slightly higher with T SP treatment. This is opposite to the common practice of material development increasing elongation usually causes decrease s in the strength. 62 Figure 4 - 1 Ultimate T ensile S trength (a) and E longation to F ailure (b) of I ngots M ade of A4047 P owders without and with TSP T reatment . (a) (b) 63 4 .3.2 As - cast Microstructure The improved mechanical properties of samples with TSP treatment can be explained by the microstructure . TSP treatment resulted in a significant refinement of both the primary Al and eutectic microconstituent in A4047 alloys. Figures 4 - 2 (a) (f) show optical micrographs of the ingots made from the A4047 powder without an d with the TSP treatment. A4047 has a near - eutectic Al - Si composition , with the primary Al phase and eutectic microconstituent. Figure 4 - 2 (a), (c), and (e) display the microstructure of the A4047 control with the objective lens of 5x, 20x and 100x, respec tively, showing Al dendrites and irregular Si flakes. With TSP treatment, under the same casting and cooling conditions, the primary Al phase was gradually refined , as shown in Figure 4 - 2 (b) and (d), and the Si flakes in the eutectic colonies were broken down to very fine fibrous phase, as shown in Figure 4 - 2 (f). The secondary dendrite arm spacing (SDAS) of A4047 ingots was gradually reduced with TSP treatment. The SDAS values of primary Al in both A4047 and TSP - treated A4047 were measured using the ImageJ software as shown in Figure 4 - 3 . The SDAS of primary Al in A4047 control was 44.9 µm while the average SDAS in TSP - treated A4047 was decreased to 15.3 µm. The reduced SDAS with TSP treatment corresponds with the improved tensile strength. In addition, the size of the eutectic Si phase in the A4047 ingots was sign ificantly reduced with TSP treatment. The values of the eutectic Si size in both A4047 and TSP - treated A4047 were measured using the ImageJ software , as shown in Figure 4 - 4 . The average grain size of eutectic Si in the A4047 control was 9.36 µm 2 while the average grain size of eutectic Si in TSP - treated A4047 was decreased to 1.01 µm 2 . Therefore, t he improved elongation of samples with TSP treatment can be explained by the modification of Si morphology from flaky to fibrous. 64 Figure 4 - 2 Optical M icrographs of Al - Si A lloys S howing the E ffect of TSP on the P rimary Al and E utectic Si M orphology of th e as - C ast I ngot: (a) T he A4047 C ontrol with an O bjective L ens of 5x, (b) T he TSP - treated A4047 with an O bjective L ens of 5x , (c) T he A4047 C ontrol with an O bjective L ens of 20 x , (d) T he TSP - treated A4047 with an O bjective L ens of 20x , (e) T he A4047 C ontrol with an O bjective L ens of 100 x , and (f) T he TSP - treated A4047 with an O bj ective L ens of 100 x . 65 * 100 dendrites per condition were measured to get an average and standard deviation values. Each dendrite with at least four arms was select. Figure 4 - 3 Effect of TSP T reatment on the S econdary D endrite A rm S pacing (SDAS) of the A4047 I ngots . 66 Figure 4 - 4 ImageJ Analysis of (a) the A4047 I ngot and (b) the TSP - treated A4047 I ngot . 67 4 .3.3 Fracture S urface Si platelets are shown in Figure 4 - 2 (e ) contribute to the brittle fracture and low ductility as the crack propagates through the interface between Si platelet s and the primary Al phase. In contrast, a continuous Al matrix with dispers ed Si cuboids, shown i n Figure 4 - 2 (f ), resulted in higher ductility when the crack propagated through the mixture of ductile primary Al and Si phases. Fractured samples made of the A4047 alloys , after tensile testing, without and with TSP were imaged are shown in Figures 4 - 5 and 4 - 6 , respectively. 68 Figure 4 - 5 A Fractured A4047 T ensile B ar: (a) Photograph of F ractured S ample, (b) Stereo P hotograph of F racture S urface, and (c) Back - scatter ed E lectron M icrograph of F racture S urface. (a) (b) (c) 69 Figure 4 - 6 A Fractured TSP - treated A4047 T ensile B ar: (a) Photograph of F ractured S ample, (b) Stereo P hotograph of F racture S urface, and (c) Back - scatter ed Electron M icrograph of F racture S urface. (a) (b) (c) 70 The subfigures (a) in Figures 4 - 5 and 4 - 6 are the photographs of fracture specimens after tensile testing from ingots without and with TSP treatment, respectively. The control sample in Figure 4 - 5 (a) had a 5% el ongation to failure and showed very little deformation at the fracture. The TSP - treat ed sample, shown in Figure 4 - 6 (a), with an 18% elongation to failure, clearly showed a necking phenomenon at the fracture point. Further indications of brittle failure of the sample without TSP treatment were evidenced by the shining surface in the stereo photograph of figure 4 - 5 (b) and large cleavage in the electron micrograph (Figure 4 - 5 (c). In comparison, failure sample with T SP treatment showed dull fracture surface in Figure 4 - 6 (b), and denser dimples in Figure 4 - 6 (c). In summary, cleavage - like brittle fractures were observed on the fracture surface of A4047 in Figure 4 - 5 (b). However, the cleavage - like facet s became smalle r or the larger cleavage - like facet s are replaced by dimples with the addition of TSP in Figure 4 - 6 (b), which makes the surface much rougher and thus results in higher elongation to failure. 4 .3.4 Solidification TSP depressed the undercooling and reduced the growth temperature during the Al - Si eutectic reaction. The cooling curves of the A4047 ingots with and without TSP treatment are shown in Figure 4 - 7 . The characteristic temperatures for the eutectic reactions obtained from these curves are shown in Ta ble 4 - 2 . As shown in Table 4 - 2 , the undercooling in the A4047 control was about 3 K , while there was no undercooling with the TSP treatment. Moreover, the eutectic temperature shown in the A4047 control was 8 48 K ( 57 5 °C ), which is the typical Al - Si eutectic temperature [ 7 ] . The TSP - treated alloy showed the eutectic temperature decreased to 83 1 K (5 58 °C). The se results contrast with Sr - modified samples , which show a decrease in eutectic temperatures with increased undercooling temperatures. 71 The increase of the undercooling with Sr modification can be explained by the restricted nucleation theory. The Sr addit ion decreases the number of Si - Si covalent bonds, reducing the number of silicon nuclei in the melt, and increas es the undercooling , due to the hindrance eutectic silicon nucleation [ 63 , 106 ] . The decrease of the eutectic growth temperature can be explained by two restric ted growth theories. One of them [ 26 , 68 ] refers to the poisoning of re - entrant edges. Sr sits at re - entrant edges, acting as a deactivator to inhibit Si growth rather than a nucleation site to promote the growth process. Instead of the Si growing along <112> directions anisotropically, the isotr opic growth behavior along multiple directions takes place in the modified Al - Si eutectic colonies. Lu and Hellawell [ 70 ] claimed that Sr is expelled into the liquid ahead of the solidification front, because of the extremely low solubility of Sr in both aluminum and silicon. Therefore, Sr elements at trace levels are absorbed onto the silicon layers, which can induce twinning by poisoning silicon growth ledges. The morphology change of Si crystals from flakes to fibers can be explained by the poisoning of the re - entrant edges and/or growth ledges to induce Si twinning. Therefore, th e increased undercooling and depressed eutectic growth temperature during solidification of Sr - modified Al - Si alloys can be explained by the presence of additional Sr . However, the undercooling of A4047 was depressed significantly with TSP incorporation . Instead of increasing t he interfacial energy between the eutectic liquid and solid , as in the case of Sr addition, TSP molecules act as heterogeneous nucleation sites during the eutectic reaction to decrease the interfacial energy during the eutectic re action, lower the energy barrier , and thus aid the nucleation of eutectic Si. The A4047 eutectic temperature was also depressed with TSP incorporation. The higher nucleation density and a flaky - fibrous transition in the silicon morphology lead s to a large solid - liquid interfacial area during solidification, resulting in the diffusion of Si atoms at the interface and requiring relatively low temperature 72 during crystal growth. In summary, on the basis of the observed refinement of both the primary Al and eute ctic Si microconstituents, the proposed modification mechanism for Al - Si alloys with TSP incorporation is that the nano - sized cage structure of TSP serves as a heterogeneous nucleation site to refine primary Al and modify eutectic Si. Figure 4 - 7 Cooling C urves of A4047 and TSP - treated A4047 Alloys . 550 560 570 580 590 600 610 0 5 10 15 20 25 30 Temperature, Degree Celsius Time, Second A4047 Control TSP-treated A4047 73 Table 4 - 2 Eutectic R eaction T emperatures for A4047 with and without TSP E xtracted from C ooling C urves . TSP - treated System T M in [K( °C )] T G [K( °C )] T [K( °C )] A4047 C ontrol 84 5 K (57 2 °C ) 8 48 K (57 5 °C ) 3 K (3 °C) TSP - treated A4047 831 K (5 58 °C ) 83 1 K (5 58 °C ) 0 K ( 0 °C ) * T Min : the minimum eutectic reaction temperature of Al - Si; T G : the eutectic growth temperature of Al - Si; T: the undercooling. Temperature difference between T Min and T G . 74 4 .3.5 Influence of T reatment T i me To investigate the reaction kinetics of A4047 powders and TSP, the effect of dwell time of A4047 powders in TSP e thanol solution was studied. The step 1 and step 3 in section 3. 2.2 were kept the same. However, step 2 was changed by varying the dwell time of A4047 powders in TSP solution. Clean A4047 p owders were soaked in TSP solutions for 4, 8, 24 hours, and 9 days. Figure 4 - 8 shows optical micrographs of ingots made of A4047 powders with different treatment time s in TSP solution. There was no change in the morphology of Si eutectic after 4 - hour trea tment over the control sample. The morphology of eutectic colonies was refined after 8 - hour treatment. 8 - hour, 24 - hour , and 9 - day treatment showed the similar refined eutectic morphology. This microstructure change suggests that the reaction between A4047 and TSP is kinetically controlled by the time. The hardness improvement is corresponding with the changes in microstruc tures. Table 4 - 3 lists the Vickers hardness value s of the ingots made of the powders with different treatment time. Still, there was no change in hardness after 4 - hour treatment over the control. Hardness improvement s after 8 - hour, 24 - hour, and 9 - day treat ment s w ere observed, which was consistent with the refine d morphology. 75 Figure 4 - 8 Optical M icrograph of an I ngot M ade of the A4047 P owders : (a) without TSP T reatment , (b) 4 - hour T reatment , (c) 8 - hour T reatment , (d) 24 - hour T reatment , and (e) 9 - day T reatment . (a) (b) (c) (d) (e) 76 Table 4 - 3 Microhardness of the TSP - treated A4047 A lloys . Samples Vickers Hardness (kg/mm 2 ) A4047 control 61.75 ± 2.5 0 4h - TSP - treated A4047 61.40 ± 1.83 8 h - TSP - treated A4047 65.93 ± 2.20 24 h - TSP - treated A4047 67.65 ± 2.93 9d - TSP - treated A4047 68.28 ± 1.17 * The applied force and dwell time were 100gf and 15 second, respectively. 10 measurements for each sample were performed to obtain the average and standard deviation values. 77 4 .3.6 Sequential D ilution 1) Indentation microhardness The microhardness of the eutectic microconstituent of the A4047 ingots improved with TSP treatment. Vickers hardness values of the ingots made from A4047 powder with and without TSP treatment are shown in Table 4 - 4 . The hardness of the A4047 ingot increased from 61.75 to 69.78 Hv with TSP treatment (about 13% improvement). These results a re consistent with tensile properties that show that the ultimate tensile strength and ductility increased to 15% and 250% with TSP treatment, respectively . Table 3 also illustrates the microhardness values of sequentially diluted TSP - treated A4047 ingots. The ingot made of 50% untreated + 50% untreated A4047 with the hardness value of 62.60 Hv showed the similar hardness property with the A4047 control while the average hardness values of sequentially diluted TSP - treated A4047 were higher than that of A404 7 control. Moreover, the hardness of the A4047 ingot with only 150 ppm TSP increased to 64.33 Hv , suggesting that 150 ppm TSP can show enhancement in the hardness of eutectic colonies of Al - Si alloys. 78 Table 4 - 4 Mi crohardness of the D iluted A4047 A lloys . Samples Vickers Hardness (kg/mm 2 ) A4047 control 61.75 ± 2.5 0 TSP - treated A4047 69.78 ± 1.83 50% untreated A4047 + 50% untreated A4047 62.60 ± 2.20 50% untreated A4047 + 50% TSP - treated A4047 69.78 ± 2.93 75% untreated A4047 + 25% TSP - treated A4047 68.60 ± 1.17 94% untreated A4047 + 6% TSP - treated A4047 64.65 ± 1.33 98.5% untreated A4047 + 1.5 % TSP - treated A4047 64.33 ± 1.96 * The applied force and dwell time were 100gf and 15 second, respectively. 10 measurements for each sample were performed to obtain the average and standard deviation values. 79 2) Microstructure T he modified microstructure of the eutectic microconstituent of TSP - treated A4047 was maintained aft e r sequential dilution. Figure 4 - 9 shows optical micrographs of sequentially diluted ingots. The control one with 50% untr eated A4047 addition in Figure 4 - 9 (a) was under the same casting process, which showed the similar morphology to untreated A4047 ingot in Figure 4 - 2 (e). However, the ingots with 50% TSP - treated A4047 and 50% untreated A4047 in Figure 4 - 9 (b) maintained the same refined morphology as the TSP - treated A4047 master alloy in Figure 4 - 2 (f). T he modified morphology for all the diluted ingots was maintained as shown in Figures 4 - 9 (b) - (e). The morphology of the eutectic colonies in A4047 with 150 ppm TSP was still refined. Therefore, the improved hardness in the sequentially dilute d TSP - treated ingots in Table 4 - 4 corresponds to the maintained morpholo gy after dilution in Figure 4 - 9 . These results also suggested that TSP can survive after multiple re - melting in 800 ° C furnace , which possessed a great potential for recycling applications. 80 Figure 4 - 9 O ptical M icrographs of D iluted A4047 I ngots: (a) 50% U ntreated A4047 + 50% U ntreated A4047, (b) 50% U ntreated A4047 + 50% TSP - treated A4047, (c) 75% U ntreated A4047 + 25% TSP - treated A4047, (d) 94% U ntreated A4047 + 6% TSP - treated A4047 . 81 3) Solidification TSP depressed the undercooling and reduced the growth temperature during Al - Si eutectic reaction , as shown in Figure 4 - 7 . More interestingly, the depressed undercooling and eutectic temperature were dependent on the TSP level. The cooling curves of A4047 after sequential dilution are shown in Figure 4 - 10 . The cooling curves of TSP - treated A4047 have been overlaid along the time - axis for clarity. The characteristic temperatures for the eutectic reactions obtained from these curves are shown in Table 4 - 5 . As shown in Table 4 - 5 , the undercooling in the A4047 control was about 3 K , while there was no undercooling with TSP treatment. Moreover, the eutectic temperature shown in the A4047 control was 850 K ( 577°C ) . The TSP - treated alloy showed the eutectic temperature decreased to 833 K (560 °C). The depressed undercooling and eutectic temperature were dependent on the TSP level in Al - Si ingots in Table 4 - 5 . The undercooling increased as the TSP level decreased. These results give a further indication that TSP reduces t he solid - liquid interfacial energy to modify eutectic Si. 82 *The cooling curves of the TSP - treated A4047 have been overlaid along the time - axis for clarity. Figure 4 - 10 Cooling C urves of D iluted TSP - treated A4047 i ngots . Table 4 - 5 Eutectic R eaction T emperatures E xtracted from C ooling C urves . TSP - treated System T M in [K( °C )] T G [K( °C )] T [K( °C )] A4047 control 846.5 K (573.5 °C ) 850.0 K (577.0 °C ) 3.5 K (3.5 °C) TSP - treated A4047 833.0 K (560.0 °C ) 833.0 K (560.0 °C ) 0 K ( 0 °C ) 75 % untreated A4047 + 25% TSP - treated A4047 839.5 K (566.5 °C ) 839.5 K (566.5 °C ) 0 K ( 0 °C ) 94% untreated A4047 + 6% TSP - treated A4047 844.5 K (571.5 °C ) 845.0 K (572.0 °C ) 0.5 K ( 0.5 °C ) 98.5% untreated A4047 + 1.5 % TSP - treated A4047 843.5 K (570.5 °C ) 846.0 K (573.0 °C ) 2.5 K ( 2.5 °C ) * T Min : the minimum eutectic reaction temperature of Al - Si; T G : the eutectic growth temperature of Al - Si; T: the undercooling. Temperature difference between T Min and T G . 83 4 .3.7 Proposed Modification Mechanism With TSP addition, the morphology of Si crystals is modified to the form of a fibrous structure , resulting in improve d ductility. The undercooling of t he Al - Si eutectic decreases with the addition of TSP and was inversely proportional to the amount of TSP in the melt. The change in undercooling during solidification can be explained by nucleation, which involves the formation of silicon nuclei. The free energy barrier for heterogeneous nucleation can be represented by Eq 4 - 1 Equation 4 - 1 In Eq. 3 - 1, is the w ett ing angle of the melt on pre - existing solid particle is the solid/liquid i nterfacial energy , is the u ndercooling , and S f is the e ntropy of fusion . Eq. 4 - 1 states that the energy barrier for nucleation decreases with decreas ing interfacial energy. Th e presence of a pre - existing solid particle will reduce the interfacial energy only if Si can easily form on particle surface. In general, the lattice mismatch between a pre - existing particle and solute have to be less than 10% to facilitate the nucleation process [ 107 ] . The Si - O - Si cage structure in TSP is in the form of cubic cluster with a diameter of 5.343 Å (the lattice parameter of Si is 5.413 Å ) , where t he Si - O - Si angles are in the range of 145 - 152 o , Si - O bond length is 1.55 - 1.65Å , and atomic radius of Si and O are 1.176 and 0.60Å . Therefore, TSP has a low lattice mismatch with the Si crystal, suggesting that the TSP can reduce the interfacial energy during Si nucleation process. Table 4 - 5 shows that addition of TSP to an Al - Si eutectic alloy is accompanied by reduc ed undercooling , which supports the fact that TSP aids nucleation of the eutectic silicon by decreasing the interfacial energy of the melt. Hence, the presence of TSP in nearly eutectic Al - Si alloys aids the nucleation of eutectic Si due to low lattice mis match with eutectic Si. 84 Growth begins after nucleation. In un - modified Al - Si eutectic alloys, t he Si crystal is bounded by 111 habit planes with two parallel twins [ 108 ] . T he 141° re - entrant corner is the most favorable site for Si crystals to nucleate, and eutectic Si grows along <112> directions at the reentrant corner , as shown in Figure 4 - 11 . This growth mechanism is well known as twin plane re - entrant edge (TPRE) [ 108 ] . What happens after adding the TSP? The addition of the TSP decrease d the eutectic growth temperature. Th e eutectic Si growth temperature was inversely proportional to the amount of TSP in the melt. Co m bin ing the solidification results and the TPRE mechanism in the unmodified Al - Si sys tem, a restrict ing TPRE growth theor y with TSP addition is proposed: TSP bonds with Al to slow down the Al segregation from Al - Si melt during eutectic reaction, leading to the microstructural refinement of Al - Si eutectic. To summarize, TSP modification of eutectic Si crystals in Al - Si alloys is associated with two mechanisms: reducing the interfacial energy to aid nucleation and slow ing down the Al segregation from Al - Si melt to suppress Si growth , which leads to morphological transformation and mechanical enhancements. 85 Figure 4 - 11 Schematic I llustrating 2D V iew of the T win P lane R e - entrant E dge G rowth M echanism . 4 .4 Conclusion s In this chapter , a new process was developed to enable the in - situ kinetic reaction of TSP and A4047. The microstructure of the eutectic microconstituent of A4047 with TSP treatment was modified, which was still maintained with only 150 pp m TSP after sequential dilution, s hedding light on the Sr fading issues. A modified microstructure was created due to the diminishment of undercooling, while the eutectic temperature changed from 850 K (577°C) to 833 K (560 °C) with TSP treatment during solidification. The change of morpho logy with TSP treatment was associated with mechanical enhancement of the A4047 alloy. These results suggest that TSP aids nucleation of eutectic Si by decreasing the interfacial energy of the melt, while inhibits Si growth by slow ing down the Al segregation from Al - Si melt , leading to morphological transformation and mechanical enhancement. 86 Chapter 5 Development of Al - Si - TSP Master Alloys 5 .1 Background TriSilanol POSS (TSP) as a novel chemical modifier was used to modify the microstructure of Al - Si based casting to increase the ductility without sacrificing the strength of the unmodified alloy. In chapter 3 , microstructure modification was obtained in a nearly eutectic alloy, A4047(Al - 12Si), after melting TSP - treated A4047 powders at laborator y scale. T he objective in this chapter is to develop a TSP master, which can be easily incorporated into base alloy in aluminum foundries at large scale . Two important goals need to be achieved: 1) the master must meet the temperature range of the thermal stability of TSP; 2 ) the master must assure that the reaction between TSP and Al occurs efficiently to refine microstructure. There are three types of A4047 - TSP master in this study: loose powder , cold - pressed puck and pre - melted ingot. At the laboratory scale, highly modified Si in A4047 after sequential dilution was obtained as shown in chapter 3 . However, the A4047 - TSP master in the form of cold - pressed puck or ingot will be more practicable to be added into the molten alloy at larger scale. The A4047 - TSP master in the form of powder requires fast - heating - rate furnace and narrow process window to allow TSP to react with Al, which is difficult to process in aluminum foundries. Therefore, two type s of A4047 - TSP master , cold - pressed puck and pre - melted ing ot, were used in this study. Besides the form of master, TSP concentration and relative amount of master alloys were investigated at 25 - kg laboratory scale. Al - 12Si and Al - 7.5Si were selected as base alloy s to study the effectiveness of the A4047 - TSP maste r. 87 5 .2 Experimental 5 .2.1 Materials and Melt Preparation Pure Al ingot consists of >99.5% Al. Al - 50Si master ingot. Chemical composition (nominal %): 50 Si, balance Al. A4047 powder - 325 mesh (44 µm or less) from Johnson Manufacturing (Princeton, IA). Chemical composition (nominal %): 12 Si, balance Al . Trisilanolphenyl - POSS (TSP) , obtained from Hybrid Plastics, Inc. (Hattiesburg, MS), was used in this study. The material s in the form of 5 - kg ingots were melted in 35 - kg capacity silicon carbide crucibles, using an electric melting furnace. In each cast, the melt temperature was maintained at 720 o C before sampling. 5 .2.2 TSP Incorporation Process 1) Cold - pressed Puck The TSP was disso lved in base alloy via puck, where th e TSP is combined with A4047 powder then pres sed to form a cylindrical puck. A puck is 24.5 mm in diameter and pressed with a force of 40~45 KN (which is about 78 ~ 88 MPa) at 200 o C and hold for 2 0 min. 2 ) Pre - melted Ingot The TSP - treated A4047 ingot as t he master alloy was added into base alloy at 720 °C to get the Al - 7.5Si ingot with about 1000 ppm level of TSP. The A4047 - TSP master alloy was prepared by melting the A4047 - TSP mixture at 720 °C . 5 .2.3 Casting 27 kg of alloy was prepared by adding pure Al and Al50Si master alloy to melt to achieve composition. The melt was degassed through 2 min injection flux followed by 7 min N 2 purge. The molten alloy was poured to a torpedo - shaped steel mold to form torpedo shape sample. The mold was maintained at 175 o C using the heated oil. 88 5 .2.4 Characteristics 1) Optical Emission Spectrometry A sample of molten alloy was poured to a spectrometer disk to perform qualitative and quantitative analysis of the elements with the SPECTROMAXx arc spark OES metal analyzer. A t least six measurements of each sample were performed. 2 ) Metallography The cast ingot was cut , exposing the middle interior region of the sample, and mounted with epoxy resins . Metallographic samples were prepared by grinding with SiC paper and then poli shing with 5/ 5 The polished mount was e 30s prior to imaging. O ptical micrographs were taken using the inverted metallographic microscope , Nikon ECLIPSE MA200, wi th the magnification up to 1,000X. 5 .3 Results and Discussion 5 .3.1 Melt Characterization The base alloy s used for this chapter w ere prepared from commercially - pure aluminum ingot with Al - 50Si master alloy addition to achieve the composition listed in Tab le 5 - 1 . The Al - 7.5Si base alloy composition is very similar to 300 series casting alloys except for Mg and Cu additions. The two base alloys contain 0.15% Fe due to it being a secondary alloy. Table 5 - 1 Chemical C omposition in wt% of the Al - 7.5Si B ase A lloy . Base Alloy Si Fe Al Al - 7.5Si 7.51 0.15 Balance Al - 12Si 12.07 0.16 Balance 89 5 .3. 2 Form of Master Table 5 - 2 summaries the alloys prepared using two types of A4047 - TSP master. Two base alloys are used in this study, Al - 12Si and Al - 7.5Si. The master forms include cold - pressed puck and pre - melted ingot. The A4047 - TSP master contains 2wt% TSP in A4047 matrix. Table 5 - 2 Naming C onventions of A ll C ast S pecimens for M aster F orm S tudy . Name Master Amount of master (wt%) Master form Al - 12Si - P A4047 25 Cold - pressed puck Al - 12 Si - TSP - P A4047 - 2TSP 25 Cold - pressed puck Al - 12Si - I A4047 25 Pre - melted ingot Al - 12 Si - TSP - I A4047 - 2TSP 25 Pre - melted ingot Al - 7.5 Si - P A4047 25 Cold - pressed puck Al - 7.5 Si - TSP - P A4047 - 2TSP 25 Cold - pressed puck Al - 7.5 Si - I A4047 25 Pre - melted ingot Al - 7.5 Si - TSP - I A4047 - 2TSP 25 Pre - melted ingot Figure 5 - 1 shows the optical micrographs of Al - 12Si castings with and without TSP via cold - pressed puck process. It can be seen that the eutectic Si exhibit well - modified structure after the addition of TSP pucks. This modified structure is similar to the TSP - treated A4047 alloy via dip coating process. The same Si morphology was observed in the Al - 12Si casting with TSP vis pre - melted ingo t process in Figure 5 - 2. In terms of Si morphology, there is no significant difference between the puck - form master and the ingot - form master. 90 Figure 5 - 1 Optical M icrogra phs of A s - cast Al - 12Si A lloys with and without TSP U sing C old - pressed P uck P rocess: (a) the Al - 12Si - P B ase A lloy with the Objective Lens of 20x, (b) the Al - 12Si - TSP - P with the Objective Lens of 20x, (c) the Al - 12Si - P B ase A lloy with the Objective Lens of 100x, and (d) the Al - 12Si - TSP - P with the Objective Lens of 100x . (a) (b) (c) (d) 91 Figure 5 - 2 Optical M icrographs of A s - cast Al - 12Si A lloys with and without TSP using P re - melted I ngot P rocess: (a) the Al - 12Si - I B ase A lloy with the Objective Lens of 20x, (b) the Al - 12Si - TSP - I with the Objective Lens of 200, (c) the Al - 12Si - I B ase A lloy with the Objective Lens of 100x, and (d) the Al - 12Si - TSP - I with the Objective Lens of 100x . (a) (b) (c) (d) 92 Table 5 - 3 SDAS of the Al - 12Si A lloys with and without TSP . Samples SDAS ( µm) Al - 12Si - P 45.0 ± 7 Al - 12Si - TSP - P 31.2 ± 5 Al - 12Si - I 44.7 ± 8 Al - 12Si - TSP - I 15.4 ± 4 * 100 dendrites per condition were measured to get an average and standard deviation values. Each dendrite with at least four arms was select. The secondary dendrite arm spacing (SDAS) of A4047 ingots was gradually reduced with TSP additions. The SDAS values of primary Al for all castings were measured using the ImageJ software as shown in Table 5 - 3. The SDAS of primary Al in Al - 12Si control s was around 45 µm while the average SDAS s in Al - 12Si - TSP via cold - pressed puck and pre - melted ingot process were decreased to 31 and 15 µm , respectively. Compared with the Al - 12Si - TSP prepared with puck master, the ingot master exhibited even smal ler arm spacing. In addition, the dwell time in hot furnace for the Al - 12Si - TSP alloy prepared using cold - pressed puck is at least one hour while pre - melted ingot only needs 15 mins to allow TSP to disperse into the melt. In other word, the pre - melted ingo t needs much less time to melt than the cold - pressed puck. Typical phases presented in Al - 7.5Si are primary Al and Al - Si eutectic, where eutectic Si appears as dark phase in Figure 5 - 3 (a,c) and 5 - 4 (a,c) . W ell - modified Si morphology was observed with TSP addition via cold - pressed puck in Figure 5 - 3 (b,d) or pre - melted ingot in Figure 5 - 4 (b,d) . In the Al - 7.5Si - TSP - P casting, it took one hour to melt the TSP pucks in 720 o C furnace. Less than 20 mins is needed to melt the pre - melted ingots to cast Al - 7.5Si - TSP - I alloys. Therefore, pre - melted ingot was chosen as TSP master to modify Al - Si alloys due to dissolving much quicker, saving energy and production time. 93 Figure 5 - 3 Optical M icrographs of A s - cast Al - 7.5Si A lloys with and without TSP U sing C old - pressed P uck P rocess: (a) the Al - 7.5Si - P B ase A lloy with the Objective Lens of 20x, (b) the Al - 7.5Si - TSP - P with the Objective Lens of 20x, (c) the Al - 7.5Si - P B ase A lloy with the Objective Lens of 100x, and (d) the Al - 7.5Si - TSP - P with the Objective Lens of 100x . (a) (b) (c) (d) 94 Figure 5 - 4 Optical M icrographs of A s - cast Al - 7.5Si A lloys with and without TSP U sing P re - melted I ngot P rocess: (a) the Al - 7.5Si - I B ase A lloy with the Objective Lens of 20x, (b) the Al - 7.5Si - TSP - I with the Objective Lens of 200x, (c) th e Al - 7.5Si - I B ase A lloy with the Objective Lens of 100x, and (d) the Al - 7.5Si - TSP - I with the Objective Lens of 100x . (a) (b) (c) (d) 95 5 .3. 3 TSP Concentration in Master Table 5 - 4 summarized the Al - 7.5Si alloy prepared with different composition of A4047 - TSP master. The TSP content in the A4047 - TSP master were 2%, 4% and 6%. The A4047 - master were prepared vis pre - melted ingot. 25% of A4047 - TSP master was added into the Al - 7.5Si melt . The entire cast specimens used in the TSP concentration study are listed in Table 5 - 4 . Table 5 - 4 Naming C onventions of A ll C ast S pecimens for TSP C oncentration S tudy . Name Master TSP content (wt%) in maste r Amount of master (wt%) Al - 7.5Si A4047 0 25 Al - 7.5Si - TSP - i A4047 - 2TSP 2 25 Al - 7.5Si - TSP - ii A4047 - 4TSP 4 25 Al - 7.5Si - TSP - iii A4047 - 6TSP 6 25 Figure 5 - 5 showed the optical micrographs of A4047 - 2TSP, A4047 - 4TSP and A4047 - 6TSP masters with two different magnifications. Similar to the microstructure of TSP - treated A4047 using dip coat ing method in chapter 4 , modified Si was observed in the A4047 - 2TSP master . However, TSP in A4047 - 4TSP and A40 47 - 6TSP appeared as a form of cluster in Figure 5 - 5(d) and (f). TSP in master were tending to agglomerate when more than 2% of TSP were added into A4047 base alloy. 6% TSP was the maximum amount of TSP that can be added into A4047 base alloy without sacrif icing the overall soundness of the master alloy. Figure 5 - 6 illustrated the optical micrographs of Al - 7.5Si and Al - 7.5Si - TSP - i castings with two different magnifications. 2% TSP addition in the A4047 - TSP master modified the eutectic Si from flaky to fibro us form in Figure 5 - 6 (b,d). Similar modified Si morphology was observed in Al - 7.5Si - TSP - ii and Al - 7.5Si - TSP - ii from Figure 5 - 7 , where the A4047 - TSP master 96 contains 4% and 6% TSP, respectively. To optimize the modification efficiency, the A4047 - 6TSP was selected as a master alloy. 97 Figure 5 - 5 Optical M icrographs of the A4047 - TSP M aster with D ifferent A mount of TSP: (a) the A4047 - 2TSP with the Objective Lens of 20x, (b) the A4047 - 2TSP with the Objective Lens of 100x, (c) the A4047 - 4TSP with the Objective Lens of 20x, (d) the A4047 - 4TSP with the Objective Lens of 100x, (e) the A4047 - 6TSP with the Objective Lens of 20x, and (f) the A4047 - 6TSP with the Objective Lens of 100x . (a) (b) (c) (d) (e) (f) 98 Figure 5 - 6 Optical M icrographs of A s - cast Al - 7.5Si A lloys with and without TSP: (a) the Al - 7.5Si B ase A lloy with the Objective Lens of 20x, (b) the Al - 7.5Si - TSP - i with the Objective Lens of 20x, (c) the Al - 7.5Si B ase A lloy with the Objective Lens of 100x, and (d) the Al - 7.5Si - TSP - i with the Objective Lens of 100x . (a) (b) (c) (d) 99 Figure 5 - 7 Optical M icrographs of A s - cast Al - 7.5Si A lloys with TSP: (a) the Al - 7.5Si - TS P - ii with the Objective Lens of 20x, (b) the Al - 7.5Si - TSP - iii with the Objective Lens of 20x, (c) the Al - 7.5Si - TSP - ii with the Objective Lens of 100x, and (d) the Al - 7.5Si - TSP - iii with the Objective Lens of 100x . (a) (b) (c) (d) 100 5 .3. 4 Amount of Master Three alloys were prepared with A4047 - TSP master via pre - melt ed ingot. The relative amount of master added into base alloy were 5%, 10% and 25 %. The microstructure of these three alloys were compared to optimize the amount of master in the base alloy . The entire cast specimens used in the master content study are li sted in Table 5 - 5 . Table 5 - 5 Naming C onventions of A ll cast S pecimens for M aster C ontent S tudy . Name Master Amount of master (wt%) Al - 7.5Si A4047 25 Al - 7.5Si - TSP - 5 A4047 - 6TSP 5 Al - 7.5Si - TSP - 10 A4047 - 6TSP 10 Al - 7.5Si - TSP - 25 A4047 - 6TSP 25 Figure 5 - 8 shows the optical micrographs of Al - 7.5Si and Al - 7.5Si - TSP - 25 castings with two different magnifications. The Al - 7.5Si control has primary Al and Al - Si eutectic microconstituents , where eutectic Si appears as dark phase in Figure 5 - 8 (a,c) . Well - modified Si morphology was observed with TSP addition using 25% A4047 - 6TSP master in Figure 5 - 8 (b,d). Similar modified Si morphology was observed in Al - 7.5Si - TSP - 10 from Figure 5 - 9 (a,c), where the relative amount of the A4047 - 6TSP master is 10%. However, the Si became coarse with only 5% A4047 - 6TSP in the Al - 7.5Si base alloy as shown in Figure 5 - 9 (b,d), suggesting that more than 5% of A4047 - 6TSP is needed to effectively modify Si in th e Al - 7.5Si alloy. 101 Figure 5 - 8 Optical M icrographs of A s - cast Al - 7.5Si alloys with TSP: (a) Al - 7.5Si with the Objective Lens of 20x, (b) Al - 7.5Si - TSP - 25 with the Objective Lens of 20x, (c) Al - 7.5Si with the Objective Lens of 100x , (d) Al - 7.5Si - TSP - 25 with the Objective Lens o f 100x . (a) (b) (c) (d) 102 Figure 5 - 9 Optical M icrographs of A s - cast Al - 7.5Si A lloys with TSP: (a) Al - 7.5Si - TSP - 10 with the Objective Lens of 20x, (b) Al - 7.5Si - TSP - 5 with the Objective Lens of 20x, (c) Al - 7.5Si - TSP - 10 with the Objective Lens of 100x, (d ) Al - 7.5Si - 6TSP - 5 with the Objective Lens of 100x . (a) (b) (c) (d) 103 5 . 5 Conclusions T he eutectic Si exhibit ed well - modified structure after the addition of both cold - pressed pucks and pre - melted ingots . The modified structure was similar to the TSP - treated A l - 12Si all oy s via dip coating process. P re - melted ingot s w ere chosen as the optimized TSP master form due to dissolving much quicker, saving energy and production time. Besides the form of master, TSP concentration and relative amount of master alloys were investiga ted at 25 - kg laboratory scale. Similar modified Si morphology was observed in the based alloy with the A4047 - TSP masters containing 2, 4 and 6% TSP. To optimize the modification efficiency, the A4047 - 6TSP was selected as a master alloy. Well - modified Si mo rphology was observed with TSP addition using both 25% and 10% A4047 - 6TSP master. However, the Si became coarse with only 5% A4047 - 6TSP in the Al - 7.5Si base alloys, suggesting that more than 5% of A4047 - 6TSP is requir ed to effectively modify Si in the Al - 7 .5Si alloy. 104 Chapter 6 Aluminum Silicon Binary Alloys Modification 6 .1 Background Al - Si based casting alloys have been used in automobile industry to improve energy efficiency by reducing vehicle bodyweight. However, as compared to steel, Al - Si casting alloys have shown relatively low strength and ductility, which limit their applications as structural components. To improve their overall mechanical, grain refiner or chemical modifier has been added into base alloy, which enhances the overall mechanical propert ies through microstructure refinement. For example, titanium (Ti)/ boron (B) can react with Al to form intermetallic compounds (IMCs) [ 72 ] , such as Al 3 Ti and TiB 2 . These IMCs serve as heterogeneous nucleation sites of primary Al, resulting in refined grain size. The addition of strontium (Sr) can poison the re - entrant corner or growing ledge of eutectic Si in order to modify the morphology of eutectic Si from flaky to fibrous form [ 62 , 70 ] . However, as the S i content is greater than 3%, the effect of the Al - Ti - B grain refiner on changing Al grain size diminishes significantly [ 39 , 73 ] , which is resulting from the formation of TiSi and TiSi 2 . Sr also suffers from an ac celerated oxidation at casting temperature [ 56 ] , leading to the loss of Si modification in Al - Si alloy after the second melt. Therefore, the issues of Ti poisoning and Sr fading after the addition of refiners/modifiers become major concerns for aluminum foundries and researchers. To avoid the issues listed above, TSP has been used as a novel chemical modifier to modify the microstructure of the Al - Si eutectic alloys in our previous study at laboratory scale [ 109 ] . This chapter focuses on increasing the ductility of Al - Si casting alloys via microstructure refinement in the as - cast condition using master alloy approach at 40 - kg scale , while maintaining the stre ngth of the unmodified alloy. 105 6 .2 Experimental 6 .2.1 Materials Commercially - pure aluminum ingot (P 0506), consists of >99.5% Al. Al - 50Si master ingot. Chemical composition (nominal %): 50 Si, balance Al. A4047 powder - 325 mesh (44 µm or less) from Johnson Manufacturing (Princeton, IA). Chemical composition (nominal %): 12 Si, balance Al . Trisilanolphenyl - POSS (TSP) , obtained from Hybrid Plastics, Inc. (Hattiesburg, MS), was used in this study. 6 .2.2 TSP Master Alloy Preparation Two stage process is used to produce the TSP master alloy, A 4047 powders with 6wt% TSP were pressed to form a cylindrical puck. The powders were pressed with a force of 40~45 KN at 180 o C and held for 10 minutes to produce a puck 31.75 millimeter in diamet er. The TSP puck was melted at 700 °C in a 3 - kg capacity graphite crucible, using an electric melting furnace to homogenize the master alloy . 6 .2.3 Casting 40 kg of the base alloy was prepared from pure Al with Al - 50Si master alloy additions in an electri c resistance furnace with a clay - graphite crucible at 720 o C. A 4 - kg A4047 - 6TSP master alloy was added into the molten base alloy and stirred for dispersion . The molten alloy was poured into a steel mold to form torpedo shape sample. The metal mold temperature was maintained at 150 o C using the heated oil. To study TSP high - temperature and remelt stability, samples were cast after 1.5 and 4 hours hold time at 720 o C. The remaining alloy was removed from the furnace. The TSP treated metal was remelt ed and an additional four samples were cast using 720 o C melt temperature and hold time of 24, 48, 72, and 192 hours. The entire cast specimens are listed in Table 6 - 1. 106 Table 6 - 1 Naming C onventions of A ll C ast S pecimens . Name Holding time, hrs Remelting Holding time after remelting, hrs Al - 7.5Si - - - Al - 7.5Si - TSP - 1.5h 1.5 No - Al - 7.5Si - TSP - 4h 4 No - Al - 7.5Si - TSP - 4h - R - 24h 4 Yes 24 Al - 7.5Si - TSP - 4h - R - 48h 4 Yes 48 Al - 7.5Si - TSP - 4h - R - 72h 4 Yes 72 Al - 7.5Si - TSP - 4h - R - 192h 4 Yes 192 107 6 .2.4 Characteristics 1) Optical Emission Spectrometry A sample of molten alloy was poured to a spectrometer disk to perform qualitative and quantitative analysis of the elements with the SPECTROMAXx arc spark OES metal analyzer. For repeatablility, three measurements of the base alloy were performed. 2) Cooli ng Curve Analysis An aluminum MeltLab system was used to conduct cooling curve analysis on all alloys. A steel cup with 150 - gram molten sample was placed on an insulated stand and a thermocouple was position in the center of the melt. Temperature versus ti me data was recorded at 10 points per second till the alloy completely solidified . For repeatablility, t hree cooling curves were collected for each alloy. 3) Differential Scanning Calorimetry A TA Discovery DSC250 was used to analyze the alloy. The as - cast alloy, around 8mg, was placed in a 40 alumina pan, and melted at 725 °C at a heating rate of 1 0°C/min and kept for 1 min; after pre - melting the sample was cooled to 450 °C at a cooling rate of 20 °C/min and then reheated up to 725 °C at a heating rate of 1 0°C/min. E ndothermic melting data w as recorded as a function of temperature . For repeatablility, three samples were analyzed per condition. 4) Metallography The cast torpedo bar was cut , exposing the mi ddle interior region of the sample, and mounted with epoxy resins . Metallographic samples were prepared by grinding with SiC paper and then polishing with 5/0.3 5 O ptical micrographs were taken using the inverted metallographic microscope , Olympus GX51, with the magnification up to 1,000X. 5) Tensile Testing 108 Cylindrical tens ile specimens were machined from the torpedo rod with a gauge diameter of 6.35 m m and gauge length of 25.4 mm, per ASTM - B557 - 15 standard for room temperature testing. 6 .3 Results and Discussion 6 .3.1 Melt Characterization Table 6 - 2 shows the average composition measured after each melt treatment. It is the challenging using a spark emission spectrometer to measure exactly TSP conce ntration in the alloy. However, the Si concentration was increased by 0.1 to 0.2 wt% for the samples containing TSP due to the Si element in TSP. Future study will be necessary to develop an alternative approach to characterize the TSP concentration in the alloy. Table 6 - 2 Chemical C omposition in wt% of the Al - 7.5Si Al loy with and without TSP . Alloy Si Fe Al Al - 7.5Si 7.51 0.15 Balance Al - 7.5Si - TSP 7.59 0.14 Balance 109 6 .3.2 Solidification and Melting 1) Cooling Curve Analysis TSP depressed the first arrest of Al - 7.5Si alloy during cooling. The cooling curves of the Al - 7.5Si alloy with and without TSP are shown in Figure 6 - 1. Two major temperature arrests are observed on the cooling curve of the Al - 7.5Si base alloy in Figure 6 - 1a. The first arrest represents the solidification of primary Al dendrites, while the second arrest represents the solidification of the Al - Si eutectic. In order to quantitatively compare the solidification behavior of thes e alloys, the first derivative was plotted in Figure 6 - 1. Selected characteristic temperatures taken from cooling curves are listed in Table 6 - 3. It can be noted that the nucleation, local minimum, and maximum temperatures of primary Al were reduced by 9, 5, and 6 °C with TSP addition. Also, the undercooling, which is equal to the difference between nucleation and local minimum, was changed from 7°C for the base alloy to 3°C for the a lloy containing TSP. T he undercooling and arrest temperature during the formation of the Al - Si eutectic was reduced with TSP additions, which can be explained by the restricted TPR E growth mechanism . 110 Figure 6 - 1 Cooling C urve and its F irst D erivative of (a) Al - 7.5Si and (b) Al - 7.5Si - TSP - 1.5h . (a) (b) 111 Table 6 - 3 Selected C haracteristic T emperatures T aken from C ooling C urves. Alloy TAl nuc (°C) TAl min (°C) TAl max (°C) TAl - Si nuc (°C) TAl - Si arr (°C) Al - 7.5Si 615 608 610 577 577 Al - 7.5Si - TSP - 1.5h 606 603 604 57 4 57 4 *Each value in the table is the average of three measurements. 2) Differential Scanning Calorimetry Analysis DSC heating curves of the Al - 7.5Si alloy with and without TSP are shown in Figure 6 - 2. Onset and peak melting temperature of the Al - Si eutectic are listed in Table 5 - 4. TSP additions did not change the onset and peak melting temperatures of the Al - Si eutec tic , suggesting that the formation of Al - Si eutectic with TSP additions are thermodynamically stable. H owever, there is a noticeable shoulder at around 578 °C observed from the alloy with TSP additions , which supports the proposed nucleation mechanism: TSP aids nucleation of eutectic Si due to low lattice mismatch . The peak melting temperature of primary Al and the total enthalpy during the melting are also listed in Table 6 - 4. TSP addition reduced the peak melting temperature of primary Al by about 5 °C, wh ich is consistent with the cooling curves results. 112 Figure 6 - 2 DSC H eating C urves of the Al - 7.5Si A lloy with and without TSP . Table 6 - 4 Selected C haracteristic T emperatures T aken from DSC H eating C urves . Alloy TAl - Si onset (°C) TAl - Si peak (°C) TAl peak (°C) Enthalpy (J/g) Al - 7.5Si 577 588 616 386 Al - 7.5Si - TSP - 1.5h 577 588 611 402 *Each value in the table is the average of three measurements. 113 6 .3.3 As - cast Microstructures Typical phases presented in Al - 7.5Si are primary Al and Al - Si eutectic, where eutectic Si appears as dark phase in optical micrographs. Figure 6 - 3 shows optical micrographs of the Al - 7.5Si alloy with and witho ut TSP additions. It is widely known that Sr modified the eutectic Si through poisoning the re - entrant corner or growing ledge. The arrest temperature of the Al - Si eutectic after Sr addition is decreased , while the undercooling is increased significantly. In our case, TSP serves the same role in modifying Si morphology as Sr in terms of microstructure. However, there might be a completely new mechanism for the Al - Si alloys with TSP addition based on the thermal analysis. Since there is a barely change in t he undercooling of the Al - Si eutectic during solidification, TSP gets involved during the eutectic reaction. There is also a small shoulder observed in the DSC melting curve near the onset melting of the Al - Si eutectic, suggesting that TSP might offer a n a iding effect on the nucleation of eutectic microconstituents during solidification. The secondary dendrite arm spacing (SDAS) of the Al - 7.5Si alloy was slightly reduced with TSP additions . The SDAS values of primary Al in both Al - 7.5Si and Al - 7.5Si - TSP were measured using the ImageJ software. The SDAS of primary Al in the Al - 7.5Si base alloy was 25 µm while the average SDAS in Al - 7.5Si - TSP - 1.5h was decreased to 18 µm. The reduced SDAS with TSP additions can be due to the decreased nucleation, growth temperatures during the solidification of primary Al. After the TSP addition, the primary Al arrest exhibited much less amount of undercooling as the base alloy suggesting that TSP can aid the n ucleation process of primary Al as well . 114 Figure 6 - 3 Optical M icrographs of the Al - 7.5Si A lloy with and without TSP : (a) Al - 7.5Si, (b) Al - 7.5Si with TSP A dditions after 1.5 - hour H old in 720 o C F urnace . (a) (b) 115 6 .3.4 Tensile Properties The tensile properties of the Al - 7.5Si alloy were improved with TSP additions. The engineering stress - strain curves of the Al - 7.5Si alloy with and without TSP additions at room tempe rature are given in Figure 6 - 4. Ductility increased from 4.5% with no TSP additions to 6% with TSP additions. The addition of TSP increased the ductility of the Al - 7.5Si alloy by 5 0% without sacrificing the strength. Ultimate tensile strength and ductility as determined from the stress - strain curves are shown in Figure 6 - 5. The increase in ductility is due to the modified Si morphology from flaky to fibrous form after TSP additions. Figure 6 - 4 Stress - strai n C urves of the Al - 7.5Si A lloy with and without TSP at R oom Te mperature. 116 Figure 6 - 5 Tensile S trength (a) and D uctility (b) of the Al - 7.5Si A lloy with and without TSP . (a) (b) 117 6 .3.5 Additive Stability 1) Effect of Holding Time Both solidification and melting behavior of the Al - 7.5Si alloy with TSP addition are maintained after 1.5 - and 4 - hours furnace hold at 720 o C. Figure 6 - 6 shows MeltLab cooling curves and Figure 5 - 7 shows the DSC heating curves of the Al - 7.5Si, Al - 7.5Si - TSP - 1.5h, and Al - 7.5Si - TSP - 4h alloys. In Figure 6 - 6, reduced undercooling and depressed arrest temperature of both proeutectic and eutectic microconstituent still appear after hold TSP in 720 o C furnace for 4 hours. Also, the Al - 7.5Si - TSP - 4h alloy still shows a small shoulder near the Al - Si eutectic onset. Typically, the level of Sr present in the melt decrease with time, which is one of the major disadvantages of Sr usage . It is reported that Sr can completely burns off after holding in the melt for 4 hours as oxidation is the primary mechanism for Sr loss. In our study, both the MeltLab cooling curve and the DSC heating curve of the Al - 7.5Si - TSP - 4h alloy exhibit the same characteristics as the Al - 7.5Si - TSP - 1.5h alloy, suggesting that no TSP fading can be observed after 4 - hour hold in 720 o C furnace. More importantly, the modified morphology and improved tensile properties are still maintained as shown in Figure 6 - 8 and 6 - 9 . 118 Figure 6 - 6 Cooling C urves of the Al - 7.5Si, Al - 7.5Si - TSP - 1.5h, and Al - 7.5Si - TSP - 4h A lloys . 119 Figure 6 - 7 DSC H eating C urves of the Al - 7.5Si, Al - 7.5Si - TSP - 1.5h, and Al - 7.5Si - TSP - 4h A lloys . 120 Figure 6 - 8 Optical M icrographs of the Al - 7.5Si and Al - 7.5Si - TSP - 4h A lloys. (a) Al - 7.5Si, (b) Al - 7.5Si W ith TSP A dditions A fter 4 - hour H old in 720 o C F urnace . (a) (b) 121 Figure 6 - 9 Tensile S trength (a) and D uctility (b) of the Al - 7.5Si, Al - 7.5Si - TSP - 1.5h, and Al - 7.5Si - TSP - 4h A lloys . (a) (b) 122 2) Effect of Re - Melting The microstructure modification and solidification behavior of the Al - 7.5Si - TSP alloys are still valid after remelting 4 - hour - hold Al - 7.5Si - TSP alloy and held in the 720 o C furnace for up to 196 hours. Figure 6 - 10 shows the cooling curves of all remelted Al - 7.5Si - TSP alloys with different furnace holding time. It can be clearly seen that remelting does not change the temperature characteristics of both pri mary Al and Al - Si eutectic, such as undercooling, and arrest temperatures. Minimal fade was observed after re - melting. In comparison, melt loses 30 - 50% of Sr on remelting without additional furnace holding [ 110 ] . Our study shows the eutectic Si in all remelted Al - 7.5Si - TSP alloy still exhibit well - modified Si morphology as shown in Figure 6 - 11. 123 Figure 6 - 10 Cooling C urves of the Al - 7.5Si, Al - 7.5Si - TSP - 4h, and Al - 7.5Si - TSP - 4h - R - X. X=24, 48, 72, and 192h. 124 Figure 6 - 11 Optical M icrographs of the Al - 7.5Si, Al - 7.5Si - TSP - 4h, and Al - 7.5Si - TSP - 4h - R - X. X=24, 48, 72, and 192h : (a) Al - 7.5Si, (b) Al - 7.5Si with TSP A dditions after 4 - hour H old in 720 o C F urnace, (c) Al - 7.5Si with TSP A dditions after R emelting 4 - hour - hold Al - 7.5Si - TSP A lloy and A dditional 24 - hour H old in 720 o C F urnace, (d) Al - 7.5Si with TSP A dditions after R emelting 4 - hour - hold Al - 7.5Si - TSP A lloy and A dditional 48 - hour H old in 720 o C F urnace, (e) Al - 7.5Si with TSP A dditions after R emelting 4 - hour - hold Al - 7.5Si - TSP A lloy and A dditional 72 - hour H old in 720 o C F urnace, (f) Al - 7.5Si with TSP A dditions after R emelting 4 - hour - hold Al - 7.5Si - TSP A lloy and A dditional 192 - hour H old in 720 o C F urnace. (a) (b) (c) (d) (e ) (f ) 125 6 .4 Conclusions In this study, a novel chemical modifier with extraordinary thermal stability was developed, which shows 5 0% improvement over the Al - 7.5Si alloy in ductility without sacrificing the strength. The microstructure of both primary Al and eutectic microconstituent s of the Al - 7.5Si alloy with TSP addition was refined , and the modified microstructure was well maintained after remelting with additional 192 hours held in the 720 °C furnace. The addition of TSP reduced the undercooling and arrest temperatures of both proeutectic and eutectic microconstituents and caused the formation of a small shoulder near the eutectic onset melting temperature. These results suggest that the presence of TSP in hypoeutec tic Al - Si alloys not only has aiding effect on nucleation of both proeutectic and eutectic microconstituents, but also restrict the growth ability of eutectic Si. TSP as a novel high - stable modifier is a further advantag eous alternative in recycling operat ions. 126 Chapter 7 Commercial Cast Aluminum Silicon Based Alloys Modification 7 .1 Background Lightweight materials, such as aluminum (Al), magnesium (Mg), and titanium (Ti) alloys, are essential for improving fuel economy in automotive industry [ 111 - 114 ] . The ability to integrate into vehicle depends on their properties and the associated manufacturing processes. The mechanical properties are directl y related to their microstructures which depend strongly both on the composition and the solidification process. A wide range of alloying elements are added into the alloy as grain refiners, solid solution strengtheners, or intermetallic compounds. The mec hanical strength can be enhanced by grain size reduction or precipitation strengthening mechanisms. For example, titanium (Ti)/ boron (B) can react with Al to form intermetallic compounds (IMCs) [ 72 ] , such as Al 3 Ti and TiB 2 . These IMCs serve as heterogeneous nucleation sites of primary Al, resulting in refined grain size. Mg and Si were added into Al to increase streng th by solid solution strengthening and precipitation hardening mechanism with the formation of Mg 2 Si IMCs. Cu was added into Al alloy to increase strength and fatigue life with the formation of Al 2 Cu IMCs as well. Therefore, 319 - type Al - Si - Cu casting alloy s are widely used in automotive industry in the production of engine blocks , cylinder head and body panel [ 115 ] . To further reduce body weight, high performance alloys were designed specifically for structural applications. High amount of Si offers very good fluidit y. High - Si containing Aural TM 2 - type alloys are well suited to the casting of thin and large structural components , such as shock towers, crash system, and space frames [ 116 ] . However, ductility is dramatically sacrificed when an alloy is strengthened with the alloying elements additions. This study focuses on the effects of TSP additions on the casting, solidification and mechanical properties of cast Al - Si based commercial alloys at 40 - kg scale . 127 The alloys used for this study are based on Aural TM 2 and Ford - W319, which are used in the production of structural components and engine parts respectively . 7 .2 Experimental 7 .2.1 Materials Commercially - pure aluminum ingot (P 0506), consists of >99.5% Al. Commercially - pure Cu rod, consists of >99.9% Cu. Commercially - pure Mg ingot, consists of>99.9% Mg. Al - 50Si master ingot. Chemical composition (nominal %): 50 Si, balance Al. Al - 10Mn master ingot. Chemical composition (nominal %): 10 Mn , balance Al. Al - 6Ti master ingot. Chemical composition (nominal %): 6 Ti, balance Al. A4047 powder - 325 mesh (44 µm or less) from Johnson Manufacturing (Princeton, IA). Chemical composition (nominal %): 12 Si, balance Al . Trisilanolphenyl - POSS (TSP) , obtained from Hybrid Pla stics, Inc. (Hattiesburg, MS), was used in this study. 7 .2.2 TSP Master Alloy Preparation Two stage process is used t o produce the TSP master alloy. A4047 powders with 6wt% TSP were pressed to form a cylindrical puck. The powders were pressed with a force of 40~45 KN at 180 o C and held for 10 minutes to produce a puck 31.75 millimeter in diameter. The TSP puck was melted at 700 °C in a 3 - kg capacity graphite crucible, using an electric melting furnace to homogenize the master alloy . 7 .2.3 Casting 40 kg of the base alloy was prepared from pure Al with master alloy additions in an electric resistance furnace with a clay - graphite c rucible at 720 o C. A 4 - kg A4047 - 6TSP master alloy was added into the molten base alloy and stirred for dispersion . The molten alloy was poured into a steel mold to form torpedo shape sample. The metal mold temperature was maintained at 150 o C using the heate d oil. 128 7 .2.4 Characteristics 1) Optical Emission Spectrometry A sample of molten alloy was poured to a spectrometer disk to perform qualitative and quantitative analysis of the elements with the SPECTROMAXx arc spark OES metal analyzer. For repeatability , three measurements of the base alloy were performed. 2) Scanning Electron Microscopy To examine the three - dimensional nature of the eutectic Si, a polished specimen was deep etched in 1wt% NaOH solution for 30 min at 333 K. A deep - etched microstructure was characterized in the as - cast condition using a JSM - 6610LV scanning electron microscopy (SEM). In addition, the Al grain size of the polished sample in the epoxy mount was evaluated using a JSM - 6610LV SEM equipped with electron backscatter diffraction (EBSD) analysis systems. The accelerating voltage used was 20 kV. A 3 - To reduce the charging effects inside the SEM from the nonconductive epoxy mount, the areas surrounding a sample were covered with conductive carbon paint. 3 ) Cooling Curve Analysis An aluminum MeltLab system was used to conduct cooling curve analysis on all alloys. A steel cup with 150 - gram molten sample was placed on an insulated stand and a thermocouple was position in the center of the melt. Tempe rature versus time data was recorded at 10 points per second till the alloy completely solidified. For repeatability, three cooling curves were collected for each alloy. 4 ) Differential Scanning Calorimetry A TA Discovery DSC250 was used to analyze the al loy. The as - cast alloy, around 8mg, 129 for 1 min; after pre - melting the sample was cooled to 450°C at a cooling rate of 1 0°C/min and then reheated up to 725°C at a heating rate of 10°C/min. Endothermic melting data was recorded as a function of temperature. For repeatability, three samples were analyzed per condition. 5 ) Metallography The cast torpedo bar was cut, exposing the middle interior region of the sample, and mounted with epoxy resins. Metallographic samples were prepared by grinding with SiC paper pension. Optical micrographs were taken using the inverted metallographic microscope, Olympus GX51, with the magnification up to 1,000X. 6 ) Heat Treatment Prior to mechanical testing, all samples were subjected to a T7 - type heat treatment. The heat treatme nt is of considerable practical importance since it has a strong effect on the ductility of the material. The procedure for this heat treatment is as follows: Step 1. Each sample was solutionized at 723 K for 90 minutes; Step 2. After solutionization the s ample was quenched in room temperature water; Step 3. The sample was then artificially aged at 498K for 90 minutes. 7) Mechanical Property Testing a. Tensile Testing Cylindrical tensile specimens were machined from the T7 - heat treated torpedo bar with a g auge diameter of 6.35 mm and gauge length of 25.4 mm, per ASTM - B557 - 15 standard for room temperature, 150 C, and 300 C testing. Five samples per alloy condition were tested and averaged to determine yield stress, ultimate tensile stress, and elongation t o failure for each alloy. b. Fatigue Testing 130 Cylindrical fatigue specimens were machined from the T7 - heat treated torpedo bar with a gauge diameter of 6.35 mm and gauge length of 25.4 mm, per ASTM - E466 standard for room temperature testing. The room tempe rature high cycle fatigue (HCF) test was performed under stress control on a servohydraulic testing machine. For each alloy there were 24 samples tested to develop the staircase (16 specimens) and in addition, the incremental stress tests (8 specimens). Th e HCF data were processed in - house using Random Fatigue Limit model. 7 .3 Results and Discussion 7 .3.1 Melting Characterization Table s 7 - 1 and 7 - 2 shows the average composition of Aural TM 2 and W319 measured after each melt treatment. Aural TM 2 base alloy contain s 11wt% Si with 0.3% Mg, 0.5wt% Mn, and 0.2 wt% Fe additions. W319 base alloy comprises 7.8 wt% Si, 3.3wt% Cu, 0.2 wt% Mg, and 0.1wt% Ti. The Sr level in both of Sr - containing alloys was about 125 pp m . 131 Table 7 - 1 Chemical C omposition in wt% of the Aura l TM 2 A lloys . Si Mg Mn Fe Sr (ppm) Al Aural TM 2 11. 12 0.31 0.54 0.20 - Balance Aural TM 2 - TSP 10.94 0. 34 0.53 0.27 - Balance Aural TM 2 - Sr 10.78 0.26 0.53 0.27 125 Balance Table 7 - 2 Chemical C omposition in wt% of the W319 A lloys . Si Cu Mg Ti Fe Sr (ppm) Al W319 7.7 3.3 0.2 0.11 0.12 - Balance W319 - TSP 8.0 3.2 0.2 0.13 0.13 - Balance W319 - Sr 7.7 3.3 0.2 0.13 0.15 125 Balance 132 7 .3.2 Casting Properties One of the casting defects occur during solidification is solidification shrinkage . Figure 7 - 1 shows the top - view images of Aural TM 2 castings: Aura TM l 2 base alloys, Aural TM 2 TSP alloys, and Aural TM 2 Sr alloys. Aural TM 2 base alloys shrink when they solidify, and the riser of torpedo - shaped castings is to supply the feed metal. Solidification cracks were observed on the top of Aural TM 2 base allo ys in Figure 7 - 1. This is due to the volumetric change during solidification. This volumetric change was significantly reduced with Sr addition. TSP addition also reduced solidification shrinkage of Aural TM 2 alloys. Similar effect on solidification shrin kage with TSP additions was observed on W319 alloys. Figure 7 - 2 shows a top - view image of W319 castings with and without TSP additions. Same amount of W319 melt was poured into steel mold. The casting with TSP additions shows less surface shrinkage as compared to the W319 base alloys. 133 Figure 7 - 1 A T op - view I mage of Aural TM 2 C astings: Au ral TM 2 Control, Aural TM 2 TSP, and Aural TM 2 Sr A lloys . 134 Figure 7 - 2 A T op - view I mage of W319 C astings: W319 C ontrol and W319 - TSP A lloys . 135 7 .3. 3 Solidification and Melting 1) Aural TM 2 Alloys Figure 7 - 3 shows the DSC endothermic and exothermic curves of Aural TM 2 alloys . The corresponding constituents, such as primary Al and Al - Si eutectic, are indicated on these curves. In the Aural TM 2 base alloy, two major temperature arrests are observed on the exothermic curve of the Aural TM 2 base alloy. The first arrest represents the solidification of primary Al dendrites, and the second arrest represents the solidification of the Al - Si eutectic . However, only one major temperature arrest appears on the endothermic curve. This is primarily due to the casting history, resulting in a n overlap of these two microconstituents. Figure 7 - 4 overl ays the DSC data plots of Aural TM 2, Aural TM 2 with TSP additions , and Aural TM 2 with Sr additions. There is no significant change observed on the endothermic curves for the alloys with and without TSP/Sr addition. However, TSP additions depress the Al - Si eutectic arrest temperature by about 8 °C , suggesting TSP restricts the TPRE growth of Si eutectic . The depressed temperature also appears on the exothermic curve of Aural TM 2 Sr. 136 Figure 7 - 3 Typical DSC E ndothermic ( T op) and E xothermic ( B ottom) C urves of Aural TM 2, C ontaining the P rimary Al and Al - Si E utectic M icroconstituent. 137 Figure 7 - 4 DSC R esults of Aural TM 2 ( B lue), Aural TM 2 - TSP(Red), and Aural TM 2 - Sr (Green) during H eating ( T op) and C ooling ( B ottom) C ycle s . Exo Down 138 2 ) W319 Alloys TSP depressed primary Al arrest and increased the Cu precipitate nucleation temperature of W319 alloy during solidification. The cooling curves of the W319 alloy with and without TSP are shown in Figure 7 - 5 . The typical constituents in W319 are p rimary Al, Al - Si euetectic, Al 2 Cu ( ), Al 4 Cu 2 Mg 8 Si 7 (Q) and Al 15 (Fe, Mn) 3 Si 6 ( - Fe). Three major temperature arrests are observed on the cooling curve of the W319 base alloy in Figure 7 - 1 ( a ) . The first arrest represents the solidification of primary Al den drites, the second arrest represents the solidification of the Al - Si eutectic , and the third arrest mainly represents the solidification of Cu IMCs which contain and Q phases . T o quantitatively compare the solidification behavior of these alloys, the fir st derivative was plotted in Figure 7 - 5 . Six characteristic temperatures are noted on the first derivative curves as follows: T Al nuc : the primary Al nucleation temperature T Al min : the primary Al temperature minimum (dT/dt=0). T Al max : the primary Al temp erature maximum (dT/dt=0). T Al - Si nuc : the Al - Si eutectic nucleation temperature. T Al - Si arr : the Al - Si eutectic arrest temperature. T sol nuc : The Cu IMCs nucleation temperature. Selected characteristic temperatures taken from cooling curves are listed in Table 7 - 3 . It can be noted that the nucleation, local minimum, and maximum temperatures of primary Al were reduced by 4 , 3 , and 3 °C with TSP addition s . Reduced Al - Si eutectic arrest temperature was observed with TSP additions. However , the Cu IMCs nucleation temperature was increased by 4 °C with TSP addition s. A higher slope of Cu IMCs in the cooling curves was observed with TSP additions, suggesting TS P might change the Cu IMCs evolution during solidification. 139 Figure 7 - 5 Cooling C urve and its Fi rst D erivative of (a) W319 and (b) W319 - TSP . (a) (b) 140 Table 7 - 3 Selected C haracteristic T emperatures T aken from C ooling C urves. Alloy TAl nuc (°C) TAl min (°C) TAl max (°C) TAl - Si nuc (°C) TAl - Si arr (°C) Tsol nuc (°C) W319 603 598 598 567 566 509 W319 - TSP 599 59 5 59 5 56 3 56 2 513 *Each value in the table is the average of three measurements. 141 Figure 7 - 6 shows the DSC heating plot of W319. The corresponding constituents, such as Q, , - Fe, Al - Si eutectic, and primary Al, are indicated on the heating curve. Figure 7 - 7 overl ays the DSC data plots of W319, W319 with TSP additions , and W319 with Sr additions. In the W319 base alloy, Cu IMCs , which contains Q and , appears as a single Q peak with a small shoulder. TSP additions exhibited two distinct peaks, suggesting two eutectic re actions are occurring at slightly different temperatures. Interestingly, TSP additions also increased the peak temperature and the ratio of and Q in W319, while Sr additions reduced the peak temperatures of both Q and without changing the - to - Q ratio. Therefore, TSP changes the Cu IMCs TSP additions also reduced the peak melting t emperature of primary Al , which is consistent with the cooling curves results. 142 Figure 7 - 6 Typical DSC E ndothermic C urve of W319. Al: The Primary Al, Si: Al - Si E : Al 2 Cu, Q: Al 4 Cu 2 Mg 8 Si 7 and - Fe: Al 15 (Fe, Mn) 3 Si 6 . 143 Figure 7 - 7 DSC R esults of W319 ( B lue), W319 - TSP(Red), and W319 - Sr (Green) during H eating C ycle. 144 7 .3. 4 As - cast Microstructures Figures 7 - 8 shows optical micrographs of the Aural TM 2 alloys in the as - cast condition. The microstructures consisted of primary - Al dendrite with Al - Si eutectic in the interdendritic regions. There are only small amounts of Mg, Mn, and Fe containing phas es in the Aural TM 2 alloys. The Si phase in the unmodified Aural 2 TM alloys has the typical coarse morphology as shown in 7 - 8(a). The Si morphology appeared to be refined with either TSP modification as shown in Figure 7 - 8(b) or Sr modification, Figure 7 - 8 (c) as compared to the unmodified microstructure. The secondary dendrite arm spacing (SDAS) of all Aural TM 2 alloys was 2 3 ± 3 µm . Figure 7 - 8 Optical Micrographs of the As - cast Condition: (a) Aural TM 2 Ba se, (b) Aural TM 2 with TSP, and (c) Aural TM 2 with Sr. To further investigate more details of the modified Si phase, each of the alloys was deep - etched and imaged using SEM as shown in Figure 7 - 9. The Si phase in both unmodified and TSP - modified Aural TM 2 alloys has a lamellar morphology with the Si in the TSP - modified alloys being more refined. However, the Sr - modified material has a Si morphology in the manner of fibrous form. 145 Figure 7 - 9 SEM Microg raphs of the Deep - etched Aural TM 2 Alloys : (a) Aural TM 2 Base , (b) Aural TM 2 with TSP, and (c) Aural TM 2 with Sr. EBSD samples were prepared for grain size determination. Micrographs for each of the alloys is shown in Figure 7 - 10. The grain size in the unm odified Aural TM 2 and TSP - modified Aural TM 2 alloys was similar. However, Sr modification significantly increased - Al grain size. Figure 7 - 10 EBSD Micrographs of the As - cast Aural TM 2 Alloys: (a) Aural TM 2 Base , (b) Aural TM 2 with TSP, and (c) Aural TM 2 with Sr. Figures 7 - 11 shows optical micrographs of the W319 alloys. Similar to Aural TM 2 alloys, TSP additions refined the eutectic Si in the W319 alloys. 146 In both Aural TM 2 and W31 9 - type alloys , the TSP additions modified the Si eutectic into - Al while Sr additions modified - Al grain size. Further evidence that TSP has aiding effect on the nucleation of eutectic microconstituents comes from the reduced undercooling. More importantly, t he arrest temperature of both Sr - modified and TSP - modified AuralTM 2 alloys was decreased; however, the TSP - modified material has a Si morphology in the manner of lamellar form. This suggests that, instead of poisoning the growing ledges of Si, TSP bonds with Al to slow down the Al segregation from Al - Si melt during eutectic reaction, leading to the microstructural refi nement of Al - Si eutectic. 147 Figure 7 - 11 Optical M icrographs of the W319 A lloy with and without TSP : (a) W319 and (b) W319 with TSP A dditions. (b) (a) 148 7 .3.5 Tensile Properties The tensile properties of Aural TM 2 alloys at room temperature and 300 o C were improved with TSP additions. The engineering stress - strain curves of the Aural TM 2 alloy with and without TSP additions at both room temperature and 300 o C are plotted in Figure 7 - 1 2 . Ductility at room temperature increased from 1 % with no TSP additions to 3.5 % with TSP additions. Ductility at 300 o C increased from 13% with no TSP additions to 26% with TSP additions. Both yield strength and ultimate tensile strength were maintained with TSP addition at room temperature and 300 o C elevated temperature. Yield strength, u ltimate tensile strength , and ductility as determined from the stress - strain curves at these two testing temperatures are shown in Figure 7 - 1 3 . The increase i n ductility is due to the modified Si morphology from flaky to fibrous form after TSP additions. 149 Figure 7 - 12 Stress - strain Curves of the Aural TM 2 Alloy with and without TSP at (a) Room Temperature and (b) 300 o C. (a) (b) 150 Figure 7 - 13 Yield Strength, Ultimate Tensile Strength, and Ductility of the Aural TM 2 Alloy with and without TSP Additions at Room Temperature and 300 o C . 151 7.3.6 High Cycle Fatigue Figure 7 - 14 S - N Curves of Heat - treated Aural TM 2 Alloys: (a) Aural TM 2 without Sr, (b) Aural TM 2 with TSP, and (c) Aural TM 2 with Sr. 152 Table 7 - 4 High Cycle Fatigue Stress of Heat - treated Aural TM 2 Alloys at Room Temperature. HCF Stress (MPa) Aural TM 2 without Sr 61 ± 10 Aural TM 2 with TSP 76 ± 10 Aural TM 2 with Sr 64 ± 9 The high cycle fatigue property of Aural TM 2 alloys at room temperature was improved with TSP additions. Figure7 - 1 4 shows S - N cu rves of all Aural TM 2 a lloys . The average HCF stress and standard deviation are listed in Table 7 - 4. The HCF stress of the unmod ified Aural TM 2 alloys was 61 MPa. The HCF stress of the Sr - modified Aural TM 2 alloys was 64 MPa. TSP additions increased the HCF stress from 61 MPa to 74 MPa. The improved fatigue stress with TSP additions resulted from the modified Si without altering Si geometry as well as the overall grain size. Figure 7 - 15 Fracture Surface of Fatigue Specimen: (a) Aural 2 without Sr, (b) Aural 2 with TSP, and (c) Aural 2 with Sr Alloys. 153 Figure 7 - 1 5 illustrates the fr acture surface of three fatigue specimen. The high magnification images illustrate the initiation site of the fatigue fracture surfaces. There was no porosity on fatigue crack initiation. In addition, No Fe or other IMCs was detected from the initiati on s ite in all three conditions. Therefore, the fatigue properties of all alloys were not influenced by porosity and represented the microstructure of the alloys. 7 .4 Conclusions Nano - structured trisilanol polyhedral oligomeric silsesquioxane (TSP) as a novel chemical mo difier was incorporated into Aural TM 2 and W319 - type alloys. TSP additions reduced solidification shrinkage during casting. Metallography showed refined Si in the Al - Si eutectic constituent and reduced Al SDAS , resulting in at least 100% improvement in ductility at both room temperature and 300 o C elevated temperature without sacrificing the strength . T SP - modified structure improved room temperature high cycle fatigue stress of Aural TM 2 alloys from 64 MPa with Sr additions to 74 MPa with TSP additions, providing solutions to reinforce aluminum alloys in body structures for the environmental and fuel economy benefits. Besides microstructural refinement, addition, suggesting TSP changes the Cu IMCs demonstrated its potential as a strengthening precipitate. 154 Chapter 8 Summary and Future Work 8 .1 Summary T risilanol polyhedral silsesquioxane (TSP) was first incorporated in Sn - based binary and ternary alloys, leading to microstructure refinement and high creep resistance. The ingot that made of TSP - treated Al - 12Si powders exhibited 250% improvement in ductility without the loss of strength as compared to the ingot made of the untreated A l - 12Si powders. Microstructure and thermal characteristics confirmed that the improved mechanical was due to the reduced undercooling and arrest temperatures for Al - Si eutectic reaction that resulted in well - modified Si morphology. The TSP - treated A4047 ingots ca n be further diluted from 10,000 ppm down to 150 ppm TSP without the loss of modified Si structure . The new TSP master alloy was produced to facilitate the practice addition of TSP to the m elt. Different approaches were performed to assess the level of T SP modification and find the optimum content of TSP master. The optimized A4047 - TSP master composition is 6%TSP in an A4047 alloy. 10% A4047 - 6TSP master w as used to modify Al - Si binary an d commercial alloys using traditional foundry process at a large 100 lbs. scale . In Al - 7.5 Si binary alloys, TSP additions increase the ductility from 3.5% with no TSP additions to 6.5% with TSP additions without sacrificing the strength. The microstructure of both primary Al and eutectic microconstituent s of the Al - 7.5 Si alloy with TSP addition was refined , and the modified microstructure was well maintained after remelting with additional 192 hours held in the 720 °C furnace. The addition of TSP reduced the undercooling and arrest temperatures of proeutectic and eutecti c microconstituents and caused the formation of a small shoulder near the eutectic onset melting temperature. These results suggest that TSP offers aiding effect on the nucleation and bonds with Al to slow down the Al segregation from Al - Si melt during eut ectic reaction , leading to morphological transformation . 155 In Aural TM 2 and W319 - type commercial alloys, TSP additions reduced the solidification shrinkage of castings. M etallography showed refined Si in the Al - Si eutectic constituent and reduced Al second ary dendrite arm spacing (SDAS), resulting in at least 10 0% improvement in ductility at both room temperature and 300 o C elevated temperature without sacrificing the strength . Besides microstructural refinement, 2 Cu) increases while Q ( Al 4 Cu 2 Mg 8 Si 7 ) decreased with TSP addition, suggesting TSP changes the Cu intermetallic demonstrated its potential as a strengthening precipitate . These findings support the proposed modification mechanism: TSP aids nucleation of eutectic Si by decreasing the interfacial energy of the melt , while bonds with Al to slow down the Al segregation from Al - Si melt during eutectic reaction , leading to morphological transformation and mechanical enhancemen ts. 8 .2 Future Work 8 .2. 1 U nderstanding modification mechanism with TSP addition. In this thesis, TSP has shown to reduce primary Al SDAS , modify eutectic Si, and alter the Cu IMCs evolution in cast aluminum. This microstructural change lead to enhanced ductility without sacrificing the strength. However, the detailed mechanism of structural refinement with TSP addition remain unclear. Future study will need to approve the proposed mechanism using high resolution TEM , i.e. the <112> growth directions of Si crystals, thus evidencing an important role of the TSP modifier in the Si growth. 8 .2.2 Effect of TSP on the porosity distribution and fatigue in cast aluminum alloys One of the mo st significant extrinsic defects in cast aluminum alloys is porosity , which has been shown to have a large impact on fatigue life . Typically, macroporosity with large interconnected pores formed in relatively poor - fed region during casting, resulting in uneven 156 density distribution at different locations. This is due to a combination of local solidification shrinkage and hydrogen precipitation. Future study on t he porosity distribution data will help character ize whether TSP changes density distribution via microstructural refinement to improve fatigue properties . 8 .2.3 Application of TSP in foundries The alloy development in this thesis were performed at 20 - kg laboratory scale. Develop trials in different p rocessing methods at larger scale (500kg), in cluding gravity casting and high pressure die casting , is vital to enable application of TSP in cast aluminum foundries . Other future works may include explor ation of TSP modification on other lightweight metals, such as magnesium and titanium alloys. 157 A PPENDIX 158 APPENDIX Effect of TSP on Crack Propagation in a Near Eutectic Al - Si Alloy A.1 Backgrou nd An important reason to have an understanding why and how cracks propagate and arrest in near eutectic Al - Si alloys is that it helps to explain the ductility improvement with the addition of nano - structured chemicals. The refined microstructure would either enhance or suppress the microcracks initiation near the crack tips and along crack edges [ 117 - 121 ] . To fully understand the facture mechanics associated with microcrack initiation, this appendix focus es on crack propagation in a near eutectic Al - Si alloy with TSP treatment. Four - point ben d specimens were deformed using a four - point bend fixture until a crack propagated from a notch. The deformed specimens were characterized with optical and electron micr oscopy as well as bending beam theory. A.2 Experimental A.2.1 Materials and sample preparation The Al - 12Si ingots with and without TSP were prepared with the A4047 powders treated with and without TSP. To treat the A4047 powders with TSP, the oxides from the powder surface were first removed by soaking the powder in 8 wt% HCl for 5 minutes at room temperature. T he clean powers were then soaked into the 1wt%TSP ethanol solutions for 24 hours to get the TSP - treated powders . Dried TSP - treated A4047 powders w ere melted at 1088 K (815°C). The molten melt was poured by gravity to a permanent graphite mold to form a rectangular ingot of 25 mm x 50 mm x 12 mm in size and cooled in air. Four - point bend specimen measuring 24mmx4.5mmx3.5mm was machined using el ectric al discharge machining from the cast ingots. A 0.5 mm notch was cut using a 0.3mm thick diamond saw as a crack 159 initiator. The sample was then ground with SiC paper and polish ed Stereo micro graphs were taken with an Olympus SZX 16 stereo microscope equipped with a SZX2 - RHS t ransmitted - light stand as shown in Figure A - 1 . A.2.2 Four - point bend test Bending was performed using a four - point flexure fixture as shown in Figure A - 2 . The specimen was supported with two anvils on the bottom . The force is applied by tightening four screws to push the specimen through two loading anvils. All support ing and loading anvils are hardened steel cylinders, each 4mm in diameter and 40mm long . The supporting and loading span lengths are 5mm and 22mm, respectively. The A4047 and A4047 - TSP specimen were bended with a n around 2mm crack by tightening 4 screws by 1 ¼ turn. Figure A - 3 shows the top and side view of a four - point bend fixture with a bended specimen. * Force is applied by tightening four screws. Figure A - 1 Cartoon of a F our - point Fl exure F ixture . 160 Figure A - 2 A F our - point B end F ixture with a B ended S pecimen : (a) S ide - vie w and (b) T op - view . A.2. 3 Microscopy (a) (b) 161 The Al - 12Si alloys with and without TSP were characterized using both optical and scanning electron microcopy. O ptical micrographs were taken using the inverted metallographic microscope , Nikon ECLIPSE MA200, with magnification up to 1,000X. SEM micrographs were taken using a MIRA3 TESCAN SEM . Secondary electron (SE) and backscattered electron (BSE) SEM images were captured to compare crack and fracture surface of A4047 with and without TSP treatment. A.3 Results and Discussion A.3.1 As - cast Microstructure A great refinement on both primary Al and eutectic Si microconstituent in A4047 alloys was observed with TSP treatment. Figure A - 3 shows stereo microg raphs of the 4 - point bend specimen . There corresponding microstructures are shown in Figure A - 4. Figures A - 4 (a) (d) are the optical micrographs of the ingots made from A4047 powder without and with TSP treatment. Figure A - 4 (a) and (b) display the micro structure of the A4047 control with the magnifications of 200x and 1000x, respectively, showing Al dendrites and irregular Si flakes. The primary Al phase was gradually refined with TSP treatment as shown in Figure A - 4 (a) and (c), and the Si flakes in the eutectic colonies were broken down to very fine fibrous phase, as shown in Figure A - 4 (d). 162 *A 0.5 mm notch was cut using a 0.3mm thick diamond saw . Figure A - 3 Macro of a F our - point B end specimen: (a) A4047 and (b) A4047 - TSP . (a) (b) 163 Figure A - 4 Opical M icrographs of A s - cast A4047 A lloys with and without TSP T reatment U sing D ip - coating P rocess: (a) the A4047 B ase A lloy with the Objective Lens of 20x, (b) the T4047 - TSP with the Objective Lens of 20x, (c) the A4047 B ase A lloy with the Objective Lens of 100x, and (d) the A4047 - TSP with the Objective Lens of 100x . (a) (b) (c) (d) 164 A.3.2 Microstructure after bending The stereo micrographs after bending were shown in Figure A - 5, where cracks propagated right from the 0.5 mm notch. Figure A - 5 Stereo M icrographs of B ended A4047 A lloys with and without TSP T reatment U sing D ip - coating P rocess: (a) the A4047 B ase A lloy with the Objective Lens of 7x, (b) the A 4047 - TSP with the Objective Lens of 7x, (c) the A4047 B ase A lloy with the Objective Lens of 10x, and (d) the A4047 - TSP with the Objective Lens of 10x . (a) (b) 165 Figure A - 6 illustrates the secondary and backscattered electron images of the A4047 alloy. The total crack length was about 1.9 mm from the notch. As can be seen in Figure A - 6(a) and (f), all the crack tips were close to the eute ctic Si, suggesting a strong resistance to the crack propagation by Si. Also, microcracks were initiated near the tip of the arrested crack, propagated and coalesced with the main crack to cause rupture. Figure A - 7 sho w s the SEM images of the A4047 - TSP all oy. The crack length was reduced to 0.78mm from the notch. The increased eutectic density was observed with TSP addition, which absorbed more energy to hinder crack growth. In addition, less amount of microcracks were observed with TSP additions. 166 Figure A - 6 Secondary ( Le ft) and B ackscattered ( R ight) E lectron I mages of the A4047 I ngots. A 1.9mm - long C rack was F ormed from the N otch. 167 Figure A - 7 Secondary ( L eft) and B ackscattered ( R ight) E lectron I mages of the TSP - treated A4047 I ngots. A 0.78mm - long crack was formed from the notch. 168 A.3. 3 Strain determination after bending T he bending strain s can be determined using classical beam bending theory shown in Figure A - 8 . As the arc length, L, is related to radius of curvature, R, through L= R , where is the angle. T he radius of curvature at any other radius is ( R+Y ) , where Y is the distance from the neutral axis . Therefore, the final length at any o ther radius is given by L=(R+Y) along the length of the beam can be expressed as - L0/L0= (R+Y) - R Where L0=R The radius of curvature for each bended specimen is determined using ImageJ software with a three - point circ ular ROI plugin. The distance from the neutral axis was measured as well. Figure A - 9 showed the results using the images from Figure A - 5. As can be seen, 2.53E - 4 strain was induced at notch root on A4047 when the crack propagated 1.9mm from the notch , whil e 3.02E - 4 strain was induced at notch root on A4047 - TSP when the crack propagated only 0.78mm . TSP incorporation hinders crack growth in near eutectic Al - Si alloys. 169 Figure A - 8 Strains I nduced during B ending of a B eam . 170 Figure A - 9 Stereo M icrographs of (a) the A4047 I ngot and (b) the TSP - treated A4047 I ngot A nalyzed by ImageJ . (a) (b) 171 A.4 Conclusions The microstructure of both primary Al and eutectic microconstituent s of the A4047 alloys made from the TSP - treated A4047 powders was refined. The crack path in a notched specimen was analyzed using a combination of stereo, electron microscopy and beam bend ing t echniques. Eutectic Si tends to act as a barrier to prevent cracks from growth, as all the crack tips were close to Si. The total crack length was reduced from 1.9 mm to 0.78 mm with TSP addition s . 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