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Grain refinement of B319 alloy using spark plasma sintered Al-Ti-C grain refiners Mok, Justin Alfred 2018

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Grain Refinement of B319 Alloy Using  Spark Plasma Sintered Al-Ti-C Grain Refiners by  Justin Alfred Mok  B.ASc., The University of British Columbia, 2018  A THESIS SUBMITTED IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF  MASTER OF APPLIED SCIENCE in The College of Graduate Studies (Mechanical Engineering)  THE UNIVERSITY OF BRITISH COLUMBIA (Okanagan)  August 2018  © Justin Alfred Mok, 2018 ii   The following individuals certify that they have read, and recommend to the College of Graduate Studies for acceptance, the thesis entitled:  Grain refinement of B319 alloy using spark plasma sintered Al-Ti-C grain refiners                     d  Submitted by Justin Mok in partial fulfillment of the requirements of the degree of Master of applied science   Dr. Lukas Bichler, School of Engineering  Supervisor, Professor  Dr. Dimitri Sediako, School of Engineering  Examiner, Professor  Dr. Kevin Golovin, School of Engineering  Examiner, Professor  Dr. Loïc Markley, School of Engineering Examiner, External   iii  Abstract The B319 aluminum alloy is a candidate material for automotive applications, where engineers and materials designers are seeking to improve vehicle efficiency through vehicle weight reduction. Although the alloy has many desirable properties, such as it’s excellent castability and moderate strength to weight ratio, continued efforts on further enhancement of the alloy’s strength via the grain refinement approach have been reported. However, in addition to the fundamental challenges related to solidification phenomena, there also remain several challenges related to the efficient dispersion of inoculating particles in the liquid alloy during casting, as well as its casting characteristics (e.g., fluidity) as a result of the grain refinement. Further, a method for fabricating effective and homogenous grain refining master alloys remains a challenge. This research focused on examining the grain refining ability of various grain refiners produced via the spark plasma sintering (SPS) powder metallurgy route. Analysis of solidification events was used to track phase evolution and its effect on the final grain size and also on the alloy’s fluidity. The mechanical properties of the as-cast material were measured via tensile tests according to ASTM B557M standard.  The results indicate that SPS enabled production of a homogenous grain refining master alloys for metal casting. Two grain refiners (Aluminum-Titanium Carbide [Al-TiC] and Aluminum-Titanium-Carbon Black [Al-TiCB]) were produced at high and low concentrations and added to the molten B319. The high concentration formulations for both Al-TiC and Al-TiCB had the greatest effect on the properties of the B319 aluminum alloy, as a result of achieving a reduction in the average grain size of 31% and 25%, respectively. Concurrently, these grain refiners increased the ultimate tensile strength by 6% and 8%, respectively. The flow length of the B319 was reduced by 15%, 31%, 29%, and 32% after addition of Al-1TiC, Al-5TiC, Al-1TiCB  and Al-5TiCB, respectively, thus suggesting the need for further work in optimizing the casting parameters to achieve optimal casting fillability.    iv  Lay Summary Metal casting is a manufacturing process that allows for the production of complex metallic parts. A challenge that metal casting foundries face is the ability to produce parts that contain uniform material structure, because due to the nature of many metal casting processes, the resulting parts often have a strong outer surface with weaker internal.  In this research, an additive was developed and studied to promote a uniform structure within a casting made with a B319 aluminum alloy. The effect of the additive on the B319 aluminum alloy`s mechanical properties, casting characteristics and micro structure was studied.   v  Preface This research thesis titled “Grain Refinement of B319 Alloy Using Spark Plasma Sintered  Al-Ti-C Grain Refiners” was written for the fulfillment of the Master of Applied Science  (MASc) degree at the University of British Columbia.  The contributions from the author and other collaborators are listed below: The Spark Plasma Sintering experiments were carried out at the SPS Facility at the University of British Columbia Okanagan (UBCO), Kelowna, BC, Canada by Justin Alfred Mok under the supervision of Anil Prasad and Dr. Lukas Bichler. All metal casting experiments were carried out at the UBCO metal casting laboratory by Justin Alfred Mok with assistance from Tyler Davis and Levi Lafortune. Metallographic surface etching was performed in the UBCO chemistry laboratory by Justin Alfred Mok with assistance from Levi Lafortune and under the supervision of Michelle Tofteland. All Scanning Electron Microscopy (SEM) was performed at the Charles Fipke SEM laboratory at UBCO by Justin Alfred Mok under the supervision of David Arkinstall. Optical Microscopy and grain size measurement was carried out by Justin Alfred Mok in the UBCO metallurgical laboratory. All mechanical testing of specimens, including tensile and hardness testing, was carried out by Justin Alfred Mok under the supervision of Alec Smith in the UBCO high-head laboratory. Manuscripts were compiled and written by Justin Alfred Mok under the supervision of Dr. Lukas Bichler.   vi  Table of Contents Abstract .......................................................................................................................................... iii Lay Summary ................................................................................................................................. iv Preface............................................................................................................................................. v Table of Contents ........................................................................................................................... vi List of Tables ................................................................................................................................. xi List of Figures ............................................................................................................................... xii List of Equations ........................................................................................................................... xv List of Abbreviations ................................................................................................................... xvi List of Symbols ........................................................................................................................... xvii List of Elements and Intermetallics ........................................................................................... xviii Acknowledgements ....................................................................................................................... xx Dedication .................................................................................................................................... xxi Chapter 1: Introduction ................................................................................................................ 1 1.1 Background ...................................................................................................................... 1 1.2 Research Objectives ......................................................................................................... 3 Chapter 2: Literature Review ...................................................................................................... 6 2.1 Aluminum and Aluminum Alloys .................................................................................... 6 2.1.1 Wrought Alloys ....................................................................................................... 6 2.1.2 Cast Alloys .............................................................................................................. 7 2.2 Temper Designations........................................................................................................ 8 2.3 Aluminum B319 Alloy ..................................................................................................... 8 2.3.1 Effect of Silicon (Si) ............................................................................................... 8 2.3.2 Effect of Copper (Cu) .............................................................................................. 9 2.3.3 Effect of Iron (Fe) ................................................................................................... 9 vii  2.3.4 Effect of Magnesium (Mg) ...................................................................................... 9 2.3.5 Effect of Manganese (Mn) ...................................................................................... 9 2.3.6 Effect of Nickel (Ni) ............................................................................................... 9 2.3.7 Effect of Titanium (Ti) .......................................................................................... 10 2.3.8 Effect of Zinc (Zn) ................................................................................................ 10 2.4 Mechanical Properties of Aluminum B319 Alloy.......................................................... 10 2.5 Casting Methods ............................................................................................................. 10 2.5.1 Sand Mold Casting ................................................................................................ 11 2.5.2 Permanent Mold Casting ....................................................................................... 11 2.6 Nucleation Processes ...................................................................................................... 11 2.6.1 Homogeneous Nucleation ..................................................................................... 13 2.6.2 Heterogeneous Nucleation .................................................................................... 13 2.7 Grain Refinement ........................................................................................................... 15 2.8 Master Alloys ................................................................................................................. 15 2.8.1 Titanium-Boride Grain Refiners ........................................................................... 15 2.8.2 Titanium-Carbide Based Grain Refiners ............................................................... 16 2.8.2.1 Titanium .................................................................................................... 16 2.8.2.2 Carbon Black ............................................................................................. 17 2.8.2.3 Titanium Carbide ....................................................................................... 18 2.9 Effect of Grain Refiners on 319 Alloys ......................................................................... 20 2.9.1 Alloy Microstructure and Morphology ................................................................. 20 2.9.2 Fluidity .................................................................................................................. 21 2.9.3 Mechanical Properties ........................................................................................... 21 2.10 Spark Plasma Sintering (SPS) ........................................................................................ 22 Chapter 3: Experimental Procedure ........................................................................................... 24 viii  3.1 Stages of Experimental Procedure ................................................................................. 24 3.2 SPS Master Alloy Development .................................................................................... 24 3.2.1 Materials ................................................................................................................ 24 3.2.1.1 Aluminum powder ..................................................................................... 24 3.2.1.2 Titanium powder ....................................................................................... 25 3.2.1.3 Carbon Black powder ................................................................................ 26 3.2.1.4 Titanium Carbide powder .......................................................................... 26 3.2.1.5 Powder Blending and Processing .............................................................. 27 3.2.1.6 SPS Supplies and Equipment .................................................................... 30 3.2.1.7 Sintering Schedule ..................................................................................... 32 3.3 Casting Experiments ...................................................................................................... 34 3.3.1 TP1 Mold ............................................................................................................... 35 3.3.1.1 Material...................................................................................................... 35 3.3.1.2 TPI Mold Casting: Process Parameters ..................................................... 37 3.3.2 Tensile Mold Casting ............................................................................................ 38 3.3.2.1 Material...................................................................................................... 38 3.3.2.2 Process Parameters .................................................................................... 38 3.3.3 Fluidity Mold Casting ........................................................................................... 39 3.3.3.1 Fluidity Mold: Material ............................................................................. 39 3.3.3.2 Fluidity Mold: Process Parameters ............................................................ 40 3.4 Material Characterization ............................................................................................... 42 3.4.1 Microstructure Analysis ........................................................................................ 42 3.4.1.1 Sample Preparation for Microscopy .......................................................... 42 3.4.1.2 Optical Characterization ............................................................................ 42 ix  3.4.1.3 Scanning Electron Microscopy (SEM) and Energy-Dispersive Spectroscopy (EDS) ................................................................................................. 43 3.4.2 Fluidity Analysis ................................................................................................... 44 3.4.3 Thermal Analysis .................................................................................................. 44 3.4.4 Mechanical Properties ........................................................................................... 46 3.4.4.1 Tensile Testing .......................................................................................... 46 3.4.4.2 Hardness .................................................................................................... 48 Chapter 4: Results and Discussion ............................................................................................ 50 4.1 Master Alloy Characterization ....................................................................................... 50 4.1.1 Densification of Master Alloy ............................................................................... 50 4.1.2 Microstructure of Master Alloy ............................................................................. 50 4.2 B319 Aluminum Alloy Microstructure Characterization............................................... 54 4.2.1 General Microstructure ......................................................................................... 54 4.2.2 B319 Aluminum Alloy Grain Size ........................................................................ 57 4.2.3 Thermal Analysis .................................................................................................. 59 4.3 Master Alloy Experiments ............................................................................................. 60 4.3.1 Thermal Analysis .................................................................................................. 61 4.3.2 Microstructure ....................................................................................................... 65 4.3.3 Grain Size .............................................................................................................. 67 4.3.4 Inoculating Particles .............................................................................................. 71 4.3.5 Fluidity Experiments ............................................................................................. 73 4.3.6 Mechanical Properties ........................................................................................... 82 4.3.6.1 Tensile Properties ...................................................................................... 82 4.3.6.2 Hardness .................................................................................................... 84 Chapter 5: Conclusions .............................................................................................................. 85 x  5.1 SPS Master Alloy Characterization................................................................................ 85 5.2 B319 Aluminum Alloy Characterization ....................................................................... 85 5.3 Master Alloy Experiments ............................................................................................. 86 Chapter 6: Future Work ............................................................................................................. 87 6.1 SPS Master Alloy ........................................................................................................... 87 6.2 Fluidity ........................................................................................................................... 87 Bibliography ................................................................................................................................. 88 Appendix ....................................................................................................................................... 95    xi  List of Tables Table 2.1: Wrought aluminum alloy designations [7] .................................................................... 7 Table 2.2: Designations for cast aluminum alloys [8] .................................................................... 7 Table 2.3: Aluminum B319.0 alloy elemental weight percentage composition [7] ....................... 8 Table 2.4: Mechanical properties of as-cast aluminum 319.0 alloy  [7] ....................................... 10 Table 2.5: Atomic parameters of titanium and aluminum [7] ...................................................... 16 Table 2.6: Methods for processing and utilizing carbon black [43] ............................................. 18 Table 3.1: List of all master alloys developed to refine B319 grain size. ..................................... 27 Table 3.2: Control additives .......................................................................................................... 27 Table 3.3: Polishing schedule for B319 alloy samples ................................................................. 42 Table 3.4: Etchants used to reveal as-cast microstructure ............................................................ 43 Table 3.5: Solidification characteristics of B319 .......................................................................... 45 Table 3.6: Parameters used for tensile testing .............................................................................. 48 Table 4.1: Density of SPS’d master alloys ................................................................................... 50 Table 4.2: Solidification parameters of unrefined B319 aluminum alloy. Limit of error ±2.2oC 59 Table 4.3: Fraction of solid results for unrefined B319 aluminum alloy. Limit of error ±2.2oC . 60 Table 4.4: Summarized solidification parameters of B319 alloy with and without master alloy addition. Limit of error ±2.2oC .................................................................................... 63   xii  List of Figures Figure 1.1: Flow chart of research scope ........................................................................................ 5 Figure 2.1: Representative cooling curve for a pure metal [31] ................................................... 12 Figure 2.2: Cooling curve [left] for a binary phase diagram [right] [70] ..................................... 12 Figure 2.3: Free energy change during homogeneous nucleation [32] ......................................... 13 Figure 2.4: Free energy comparison for homogeneous (blue) and heterogeneous (red) nucleation [33] ............................................................................................................................. 14 Figure 2.5: Surface energies during heterogeneous nucleation [12] ............................................. 14 Figure 2.6: Aluminum-titanium binary phase diagram  [14] ........................................................ 17 Figure 2.7: As-cast Aluminum B319: (A) etched macrostructure, (B) etched microstructure, (C) SEM of intermetallics ................................................................................................ 20 Figure 3.1: Aluminum powder: [A] macro image, [B] SEM micrograph .................................... 25 Figure 3.2: Titanium powder ........................................................................................................ 25 Figure 3.3: Acetylene carbon black: [A] macro image, [B] SEM micrograph (50,000x) ............ 26 Figure 3.4: Titanium carbide powder (CAS: 12070-08-5) ........................................................... 26 Figure 3.5: Analytical microbalance ............................................................................................. 28 Figure 3.6: Powder blend after manual shaking ........................................................................... 28 Figure 3.7: Ball mill canister and zirconia balls ........................................................................... 29 Figure 3.8: Fritsch ball mill .......................................................................................................... 29 Figure 3.9: ϕ60mm x 10mm SPS master alloy ............................................................................. 30 Figure 3.10: ϕ20mm SPS master alloy (left:Al-5TiC, right Al-5TiCB) ....................................... 31 Figure 3.11: ϕ20mm  punches, die, and graphoil for SPS ............................................................ 31 Figure 3.12: Thermal technologies 10-3 SPS machine ................................................................. 32 Figure 3.13 Thermal and pressure profile of sintering schedule used to produce SPS master alloys ........................................................................................................................ 33 Figure 3.14: Melting furnace (euclid R-85) .................................................................................. 34 Figure 3.15: Preheat furnace (euclid CF 510) ............................................................................... 34 Figure 3.16: TP1 mold set up [all units in mm] ............................................................................ 36 Figure 3.17: TP1 mold  assembly wrapped with heating tape and insulted with FiberFrax blanket ................................................................................................................................. 36 Figure 3.18: 3D CAD of fluidity mold ......................................................................................... 39 xiii  Figure 3.19: Stopper and medium thermocouple .......................................................................... 40 Figure 3.20: SEM/EDS lab station................................................................................................ 43 Figure 3.21: Spiral fluidity casting on polar graph paper ............................................................. 44 Figure 3.22: Solidification characteristics represented on a cooling curve .................................. 46 Figure 3.23:Tensile sample extraction location ............................................................................ 46 Figure 3.24: Tensile sample dimensions (dimensions in “mm”) .................................................. 47 Figure 3.25: Instron tensile machine ............................................................................................. 47 Figure 3.26: Rockwell hardness testing machine ......................................................................... 48 Figure 4.1: SEM image of Al-1TiC master alloy ......................................................................... 51 Figure 4.2: TiC particle distribution in Al-1TiC (A) and Al-5TiC (B) ........................................ 52 Figure 4.3:TiC particle distribution in Al-1TiCB (A) and Al-5TiCB (B) .................................... 53 Figure 4.4:Microstructure of unrefined B319 alloy ...................................................................... 54 Figure 4.5: SEM/EDS image of secondary phases in B319 microstructure (A) SEM image; (B) Al; (C) Fe; (D) Mn; (E) Si ......................................................................................... 55 Figure 4.6: EDS chemical analysis of the block-like CuAl2 ......................................................... 56 Figure 4.7: Grain structure of unrefined B319.............................................................................. 57 Figure 4.8: Grain size histogram for unrefined B319 alloy .......................................................... 58 Figure 4.9:Cooling curve and first derivative curve for unrefined B319 aluminum alloy ........... 59 Figure 4.10: Fraction of solid curve for unrefined B319 aluminum alloy .................................... 60 Figure 4.11: Cooling curve and first derivative; [A] B319 with Al-5TiC and Al-5TiCB addition and; [B]  B319 with Al-1TiC and     Al-1TiCB addition ......................................... 61 Figure 4.12: Fraction of solid of B319 with SPS master alloy addition ....................................... 64 Figure 4.13: SEM image of B319 with Al-1TiC [A] and  Al-5TiC [B] (X500) ........................... 65 Figure 4.14: SEM image of B319 with pure Al addition .............................................................. 66 Figure 4.15: SEM image of B319 with Al-1TiCB [A] and Al-5TiCB [B] ................................... 67 Figure 4.16: Comparing grain size between all TP1 castings. (A) Unrefined; (B)Pure Al; (C) Powder TiC: (D) Al-1TiCB; (E) Al-1TiC; (F) Al-5TiCB; (G) Al-5TiC ................. 68 Figure 4.17: Average grain size measurements of all casting experiments .................................. 69 Figure 4.18: Ti particles in the fluidity mold castings .................................................................. 71 Figure 4.19: Line scan data of Ti particle within B319 alloy after Al-5TiCB addition to fluidity mold casting [A]SEM image; [B] Ti line scan; [C] Al line scan; [D] Si line scan . 72 xiv  Figure 4.20:Temperature of liquid metal reservoir in the fluidity mold ....................................... 73 Figure 4.21: Spiral flow length of unrefined castings [A] Trial#1, [B] Trial#2 ........................... 74 Figure 4.22: Cooling curves for unrefined B319 in spiral mold ................................................... 74 Figure 4.23: Grain structure of unrefined B319 sectioned at 102mm from spiral ........................ 75 Figure 4.24: Microstructure (500x) of unrefined B319 sectioned at 102mm from spiral ............ 75 Figure 4.25: Grain structure of unrefined B319 sectioned at tip of spiral length ......................... 76 Figure 4.26: Microstructure of unrefined B319 sectioned tip of spiral length ............................. 76 Figure 4.27: Macrostructure of B319 alloy at spiral tip. [A] unrefined;  [B] Al-1TiC addition;  [C] Al-5TiC addition; [D] Al-1TiCB addition; [E] Al-5TiCB addition. ................. 77 Figure 4.28: Microstructure (2000x) of B319 alloy. [A] unrefined;  [B] Al-1TiC addition;  [C] Al-1TiCB addition; [D] Al-5TiC addition; [E] Al-5TiCB addition. ....................... 78 Figure 4.29: Flow length of all fluidity B319 experiments .......................................................... 79 Figure 4.30: SEM-EDS image of Ti cluster in B319 casting with Al-5TiCB addition. [A] SEM/EDS image; [B] Al; [C] Si; [D]; Ti; [E] Fe 81 Figure 4.31: UTS and elongation of refined and unrefined B319 ................................................ 82 Figure 4.32: Yield strength and Young’s modulus of refined and unrefined B319 ..................... 83 Figure 4.33: Hardness of refined and unrefined B319 .................................................................. 84    xv  List of Equations Equation 2.1: Interfacial energies for heterogeneous nucleation [12] .......................................... 14 Equation 2.2: Hall-Petch equation [17] ........................................................................................ 21    xvi  List of Abbreviations AA Aluminum Association ASTM American Society for Testing and Materials CAS Chemical Abstracts Service DC Direct Current DTA Differential Thermal Analysis FCC Face Center Cubic GHG Greenhouse Gases GW Global warming ID Inner Diameter OD Outer Diameter SDAS Secondary Dendrite Arm Spacing SEM Scanning Electron Microscope SPS Spark Plasma Sintering TA Thermal Analysis UBCO University of British Columbia Okanagan US United States USA United States of America UTS Ultimate Tensile Strength YS Yield Strength      xvii  List of Symbols HB Brinell Hardness C Celsius cm Centimeter g Gram Hz Hertz A Interfacial Area K Kelvin kg Kilogram km Kilometer MPa MegaPascal mm Millimeter E Modulus of Elasticity ppm Parts Per Million σ Stress T6 T6 Heat treatment HV Vickers Hardness wt% Weight %    xviii  List of Elements and Intermetallics  Al Aluminum Ti Titanium TiC Titanium Carbide B Boron C Carbon CB Carbon Black Al-Ti-B Aluminum-Titanium-Boron Master Alloy Al-Ti-C Aluminum-Titanium-Carbon Master Alloy Al-TiC Aluminum-Titanium Carbide Master Alloy Al-TiCB Aluminum-Titanium-Carbon Black Master Alloy Al-Si Aluminum-Silicon Cu Copper Zn Zinc Fe Iron Mn Manganese Mg Magnesium Al-Cu Aluminum -Copper α-Al Proeutectic Aluminum Na Sodium Sr Strontium CuAl2 Copper Aluminide xix  Al5FeSi β-Iron Intermetallic Al15(Fe,Mn)3Si2 α-Iron Intermetallic Al8FeMg3Si6 II-Iron Intermetallic TiB2 Titanium Diboride TiAl3 Titanium Aluminide (Al,Ti)B2 Titanium-Aluminum Boride Zr Zirconium Ti5Si3 Titanium Silicide Al4C3 Aluminum Carbide Ti3AlC5 Titanium-Aluminum Carbide    xx  Acknowledgements I undoubtedly need to send my most sincere gratitude to my supervisor, Dr. Lukas Bichler. He has unconditionally been able to support, teach, and guide me through my Master’s studies. His calm and optimistic approach constructed an environment of learning that was both motivational and enjoyable. Thank you for introducing me to the world of research academia and challenging me to pursue research in a field that is new and fascinating to me. It comes to no surprise that you have an impressive team of researchers working with you; I am very thankful to have been apart of your great team.  Throughout the course of this research project, I absolutely needed the help and assistance of my colleagues and peers to carry forward. Thank you Mr. Anil Prasad, for your help with all of my SPS experiments. I am so thankful for your help and your input – you are truly a great friend. I also need to extend a very big thank you to Mr. Levi Lafortune for all of his help with my metal casting experiments and metallographic surface etching. Levi’s commitment to productivity allowed me to work efficiently. From this, I was able to complete my experiments on schedule and further the progress on my research. I must also extend a big thank you to Mr. Tyler Davis for his assistance on many of my metal casting experiments and subsequent analysis.  An extended thank you must also be sent to Ms. Somi Doja, Mr. Sebastian Lemus Fonseca, Mr. Matthew Smith, and Mr. Kris Mackowiak for all of your help along the way. During the time of my broken left clavicle, you all without question volunteered to help me in any way possible – Thank you.  Last but certainly not least, I must thank Mr. Durwin Bossy and Mr. Raymond Seida for your assistance and knowledge in the machine shop. Your help allowed me to manufacture tooling and casting molds necessary for this research. Thank you, Mr. Alec Smith for your help with all my mechanical testing. Thank you, Ms. Michelle Toftland for her knowledge of chemical preparation and safety during metallographic surface preparation and Mr. David Arkinstall for your expert help on the Scanning Electron Microscope.    xxi  Dedication  To my Family and Friends who have supported me through all of my endeavours 1  Chapter 1: Introduction  This chapter provides a general overview of the present research thesis. The motivation to use aluminum alloys in engineering applications is discussed, followed by a broad overview of advancements in the strengthening of aluminum alloys. 1.1 Background Global warming (GW) has become a rising concern in recent years due to the rapidly receding polar icecaps and rising average global temperatures [1]. The cause of global warming is often attributed to the increase in the greenhouse gas emission [2]. Accumulation of greenhouse gases (GHG) such as carbon dioxide, methane, and nitrous oxide in the atmosphere further accelerates GW. In 2014, the US environmental protection agency estimated that the USA produced 6,870 million metric tons of CO2 from fossil fuels, with the transportation sector being responsible for 26% of these emission [2]. Fuel economy is used in the automotive industry to define a vehicles performance. There are many factors that affect a vehicle’s fuel economy, such as its engine size, aerodynamics, and curb weight [3]. For example, it is estimated that a 10% reduction in curb weight can reduce the fuel consumption by 6%-8%  [4]. The potent effect of curb weight on the fuel economy is often drives the automotive industry to continue to develop light-weight vehicles. One way to reduce the curb weight is to replace heavy ferrous materials with lighter aluminum alloys that posses the required strength for a given application.  Materials like aluminum have been integrated into automotive designs in an attempt to lower the curb weight [5], [6]. Next to iron (Fe), aluminum is the most abundant metal on earth [7]. Pure aluminum is produced using the Hall-Heroult process, where alumina is refined from a bauxite mineral [7]. Aluminum alloys are melted and solidified into ingots that are then sent to fabricators for further processing. Pure aluminum is typically alloyed with other elements such as magnesium, copper, and silicon to achieve specific material characteristics and properties [7], [8]. Aluminum alloys are divided into two categories - wrought and cast. Wrought alloys undergo a plastic deformation during production, which strengthens the material by the means of microstructural changes [7]. Conversely, cast alloys are melted and solidified with no plastic deformation. Aluminum alloys are relatively amenable to casting processes due to their low melting point, good fluidity, and chemical stability [7], [9]. The B319 aluminum alloy is a cast alloy often used to produce light-weight engine blocks and cylinder heads in the automotive industry [9]. By weight percent, the B319 alloy consists of 5.5-6.5% silicon, 1.2% iron, 3-4% 2  copper, 0.8% manganese, 0.1-0.5% magnesium, 0.5% nickel, 1% zinc, 0.25% titanium, and the balance of aluminum [7]. The B319 alloy has several advantages, such as a good corrosion resistance, wear resistance and moderate strength [7], [8], [10]. From a manufacturing standpoint, the alloy has excellent castability to produce near net shape parts [7], [9] . This characteristic has economic benefits for manufacturers as it can reduce the amount of wasted material and eliminate some intermediate processes required to produce the final part. Although aluminum alloys have been used for decades, there remain opportunities to further improve their properties.  The kinetics of solidification while an alloy transforms from a liquid to a solid phase play a key role on the as-cast mechanical properties of the alloy [7], [8], [11], [12]. The nature of aluminum alloy casting requires careful considerations to parameters that, directly and indirectly, effect the solidification behavior of the metal. These parameters include the pouring temperature, elemental composition, mold preheat temperature and the mold design [7], [12], [13]. The size of the grains in the aluminum alloy is determined during the alloy’s solidification [6], [8]. Faster cooling yields smaller grains, while a slower cooling yields larger grains and a weaker as-cast material [8], [13], [14]. Naturally, industrial castings with a complex shape do not facilitate uniform and homogeneous cooling throughout the casting volume, since the metal-mold interface experiences a significantly higher heat extraction than the center of the casting. Moreover, available nucleation sites are often not scattered homogenously throughout the material during solidification. As a result, the grain structure within a casting is not uniform, which makes it difficult to predict the performance of cast parts [8].  In the 1930’s, grain refinement of aluminum alloys was introduced to treat issues related to inhomogeneous grain structure [13]. Grain refinement is a process wherein the grain structure is controlled in order to achieve small grains throughout an entire casting [6], [15]–[17]. Grain refinement can be done by rapid cooling of the liquid metal during solidification or through the addition of inoculating particles. The use of inoculating particles promotes heterogeneous nucleation, which effectively reduces the average grain size [12], [13]. Heterogeneous nucleation may occur on a variety of foreign surfaces such as inclusions, oxides, mold walls, and grain refining particles within a liquid metal [12].  Titanium, Boron, and Carbon are the most common elements used as grain refiners for aluminum alloys due to their stability and nucleating capabilities [13], [18], [19]. To facilitate grain refinement, a known mass of a desired grain 3  refining element is mixed with a base material to create a master alloy. Master alloys serve as a way to easily quantize the amount of grain refiner and act as an effective transport mechanism to introduce the inoculants into the liquid metal during casting. This approach has simplicity and economical benefits during industrial practice [19]. However, there remain challenges and difficulties when using grain refiners. Fading effects from extended holding times of the particles within the melt can directly degrade the potency of the inoculants and the efficiency of the master alloys to effectively achieve a small grain size in the as-cast alloy [20]. Moreover, agglomeration of the grain refining particles within the casting can result in stress concentration sites when a casting is subjected to a tensile loading [18].   In the present work, an Aluminum-Titanium Carbide and Aluminum-Titanium- Carbon Black grain refiners were studied and their effect on the B319 alloy was investigated. The master alloys have been prepared using a Spark Plasma Sintering (SPS) powder metallurgy process, and their effect on the B319 alloy’s microstructure, mechanical properties, and casting characteristics were quantified.  1.2 Research Objectives The objective of the proposed research was to explore the feasibility of producing an effective grain refining master alloy using the spark plasma sintering process. The scope of the research thesis can be summarized as follows, while Figure 1.1 provides a graphical representation: 1) Master alloy fabrication  a. Sintering i. Aluminum-Titanium Carbide (Al-TiC) master alloy ii. Aluminum-Titanium-Carbon Black (Al-TiCB) master alloy b. Castings i. Unrefined B319 alloy ii. TiC Powder addition iii. Pure aluminum SPS pellet addition iv. Al-TiC SPS pellet addition 1. High concentration 2. Low concentration 4  v. Al-TiCB SPS pellet addition 1. High concentration 2. Low concentration 2) Experiments and Characterization a. TP1 Mold i. Microstructure ii. Thermal Analysis (TA) b. Fluidity Mold i. Microstructure ii. Thermal Analysis (TA) iii. Flow Length c. Tensile Mold i. Tensile Testing (ASTM B557M)  The thesis document has been organized as follows:  Chapter 2 provides a focused literature review pertinent to the grain refinement and microstructure characterization of B319 aluminum alloy. Chapter 3 provides a detailed description of experimental methods, equipment and analysis techniques employed in this research. Chapter 4 provides a critical discussion of the generated results, while Chapter 5 provides conclusions based on the results, as well as suggestions for future work.  5    Figure 1.1: Flow chart of research scope   6  Chapter 2: Literature Review  This section provides a concise review of the B319 aluminum alloy and the past research on the grain refinement of B319 using diverse inoculation techniques. The alloy’s characteristics, properties, and processing routes are also summarized. 2.1 Aluminum and Aluminum Alloys Aluminum and its alloys have many desirable properties, which make them amenable to many engineering applications [7]. The ability to modify the alloy composition has broadened the range of applications to accommodate many engineering design constraints and objectives. Moreover, it has allowed engineers to completely transform and optimize existing engineering designs, which were created with traditional materials and their properties in mind.  For example, aluminum alloys with silicon and magnesium offer low density, moderate strength, and good manufacturability, which make such alloys highly desirable for the production of monocoque aircraft frames. These frames enable reduced aircraft weight and increased cargo and passenger capacity [21]. Similarly, aluminum alloys with silicon and copper exhibit excellent castability that allows automotive manufacturers to produce light-weight engine blocks and castings with complex geometries [9]. Depending on the application, most raw aluminum alloys come either as a wrought or cast materials.  2.1.1 Wrought Alloys Designated with a four-digit number, wrought alloy ingots or bars  are manufactured with a hot or cold working process that strain hardens the material by means of mechanical deformation [7]. The first number in the alloy name is used to identify the primary (or principal) alloying element [7]. The second number is used to identify the modification designation for that specific alloy. The remaining two numbers are arbitrarily used to give the alloy an identification [7]. Table 2.1 below lists the seven wrought alloy classes, which are commonly used in industrial practice [7].     7  Wrought Alloy Designation Principal Alloying Element Applications 1XXX None (pure aluminum) Electrical, Thin foil 2XXX Copper  Aircraft frames 3XXX Manganese Automotive panels 4XXX Silicon Automotive welding/ brazing 5XXX Magnesium Pressure vessels 6XXX Magnesium and Silicon Irrigation piping 7XXX Zinc Bicycle frames Table 2.1: Wrought aluminum alloy designations [7] 2.1.2 Cast Alloys Cast alloys are used for metal casting applications, where the alloy is liquefied and subsequently poured into an empty cavity (i.e., a mold). The alloy undergoes solidification within the mold and the final casting is a replica of the mold cavity. The alloys often have a good fluidity and have a relatively low melting points. With many molds enabling a high heat transfer, cast aluminum alloys often enable short casting cycles and thus enable mass production of cast parts [7]. The Aluminum Association (AA) has designated cast alloys with a three-digit number, where the first number indicates the primary principal alloying element and the remaining two digits are given to identify an alloy’s unique composition [8]. Table 2.2 lists commonly available industrial cast alloys [8]. The alloy’s three-digit alloy designation can be followed by a decimal value “.0” to identify the alloy as a casting, or by “.1” to identify it as an ingot, or “.2” to identify it as an ingot with tighter elemental precision [8]. If an alloy has the same elemental composition for primary elements, but has different elemental limits on secondary elements, the alloy will be given a letter prefix before the three-digit designation.  Cast Alloy Designation Principal Alloying Element Applications 1XX.X None (pure aluminum) Experimental standards 2XX.X Copper Aircraft mounting components 3XX.X Silicon and Copper Automotive engine blocks 4XX.X Silicon Low pressure vessels 5XX.X Magnesium Cooking utensils 7XX.X Zinc Garden tools Table 2.2: Designations for cast aluminum alloys [8] 8  2.2 Temper Designations Heat treatment (or temper) of Aluminum is typically performed to homogenize the alloy’s microstructure and/or tailor the mechanical properties of the final casting to fit the requirements of a design [8], [22]. Of the cast alloy series listed in Table 2.2, the 2XX.0, 3XX.0, and 7XX.0 series cast alloys are deemed heat treatable due to proven effects of heat treatment on the alloy’s strength and hardness [8].  According to the aluminum association, the temper designation follows the alloy name, and can be one of the following six common thermal treatments used for aluminum alloys [8]: • F: as-cast • O: annealed  • T4: solution treated and aged • T5: precipitation hardened • T6: solution heat treated, quenched, and precipitation hardened  • T7: solution heat treated, quenched and overaged 2.3 Aluminum B319 Alloy The aluminum alloy B319, with its composition listed in Table 2.3, is a silicon-copper alloy from the 300-series family of casting alloys. It has a good fluidity during casting and a moderate strength at elevated temperatures [7]–[9].    Table 2.3: Aluminum B319.0 alloy elemental weight percentage composition [7] 2.3.1 Effect of Silicon (Si) Si is used to promote fluidity of the molten material during casting [7], [8], [23]. It is well documented that Al-Si alloys have the primary Al phase growing in a dendritic morphology. Si forms as a needle-like phase between the dendrite arms, and at high Si content the mechanical properties were seen to increase with increasing Si concentration [24]. As shown in Table 2.3, the B319 alloy contains between 5.5 - 6.5 wt% Si, which gives it a high fluidity without significantly compromising the alloy’s ductility. High fluidity is preferred for castings with thin sections to avoid casting defects, such as misruns and hot tears [7], [8], [25].  Alloy Si Cu Fe Mg Mn Ni Zn Ti Other Al B319.0 5.5-6.5 3.0-4.0 1.2 0.1-0.5 0.8 0.5 1.0 0.25 0.5 Balance 9  2.3.2 Effect of Copper (Cu) Cu enhances the strength and hardness of the alloy, and also allows it to become heat treatable (for compositions with 3.0-4.0%Cu), thus further improving its strength and hardness [7], [9].  Copper additions lead to the formation of a copper intermetallic compound CuAl2 and eutectic-like phases. In the presence of Si, the blocky intermetallic compound is more common. Contrary to the effect of silicon, copper as an alloying element decreases the fluidity of aluminum during casting [7], [26]. 2.3.3 Effect of Iron (Fe) Due to dissimilar crystal lattice structures, the low solubility of iron in aluminum results in the formation of intermetallic compounds, such as the FeAl3 and FeMnAl6 [7], [12], [27]. Iron helps to reduce hot tearing during solidification and improve the strength of the final solid. However, if the Fe levels exceed target values, the fluidity is compromised, since Fe reacts with Mn to create a sludge during casting and thus retards the fluid flow [8].  2.3.4 Effect of Magnesium (Mg) Magnesium is added to enhance the strength and hardness of the B319 aluminum alloy [7], [25]. Mg and Si readily form an intermetallic compound (Mg2Si) which contributes to the hardening of the alloy during heat treatment. Much like iron, solubility limits for Mg must be considered to avoid adverse affects such as matrix softening and porosity from incipient melting  [7], [25].  2.3.5 Effect of Manganese (Mn) Manganese has no direct benefit for cast aluminum alloys and is commonly considered an impurity. Thus, Mn levels must be maintained below target concentrations [7], [8].  2.3.6 Effect of Nickel (Ni) Alloying with nickel is typically done in tandem with copper to enhance the high temperature strength and hardness of the alloy [7]. Moreover, Ni reduces the coefficient of thermal expansion of the final solid and thus helps maintain dimensional tolerance of the final casting [7].  10  2.3.7 Effect of Titanium (Ti) Titanium is typically added to the alloy to refine the grains and improve the strength and hardness of the final casting. Refinement is achieved by heterogeneous nucleation [7], [8], [12], [28].  2.3.8 Effect of Zinc (Zn) Much like nickel, zinc works in tandem with magnesium and copper to make the alloy amenable to heat treatment and natural aging [7], [8]. 2.4 Mechanical Properties of Aluminum B319 Alloy The mechanical properties of the B319 alloy make it relevant to diverse industrial applications. The Aluminum 319.0 alloy’s properties (Table 2.4) have been organized into respective casting processes and temper designations; namely, sand and permanent molds and as-cast, T5 and T6 treatments, respectively. The combination of these molding processes and tempers are commonly used in industrial practice due to their proven results.  Mold & Temper Ultimate Tensile Strength (MPa) 0.2% Yield Strength (MPa) Elongation (%) Hardness (HB) Modulus of Elasticity (GPa) Sand Cast      F 185 125 2 70 74 T5 205 180 2 80 74 T6 250 165 2 80 74 Permanent Mold Cast      F 235 130 3 85 74 T6 275 185 3 95 74 Table 2.4: Mechanical properties of as-cast aluminum 319.0 alloy  [7]  2.5 Casting Methods There are many available casting methods for aluminum alloys. The selection of the most suitable casting method depends on the desired mechanical and physical properties of the as-cast material, post process operations, service conditions and manufacturing economics [8].   11  2.5.1 Sand Mold Casting Sand-mold casting, or simply sand casting, involves imprinting a replica of the desired shape to be cast into a sand and clay mold. A new mold is required for each casting. The compressive strength and permeability of the sand mold are important parameters [8], since the compressive strength must be sufficient to avoid mold collapse during mold filling, while the permeability has an effect on the casting’s surface finish and cooling behavior [11], [13], [29]. Relative to other casting methods, sand casting imposes slower cooling rates on the alloy, thus yielding larger grains and a lower strength alloy [29]. Additionally, slower cooling rates may allow hydrogen gas pick up, thus resulting in the formation of microporosity and microsgregation of solutes within the alloy [29], [30]. These inherent defects can be reduced with the proper use of permanent mold castings.  2.5.2 Permanent Mold Casting The permanent mold casting process utilizes a steel mold to make a casting. The mold can be reused to produce high volumes of identical parts. The simplest permanent mold castings are typically gravity-fed steel molds, where the hydrostatic pressure head provides the fluid the energy required to flow into the mold cavity. The benefit of the permanent mold casting process comes from the ability to reuse the mold, which effectively lowers manufacturing costs when producing high volume of parts. Additionally, the use of the steel molds enables control over the surface finish of the casting, which allows for the fabrication of high precision parts, thus requiring less post-processing. Depending on the mold material used, permanent molds can facilitate rapid heat extraction during solidification. By increasing the heat transfer and the cooling of the casting, permanent mold castings often have a fine grain size, good mechanical properties, low shrinkage porosity, and fewer gas porosity defects [8]. These improvements in properties are the result of the effect of the cooling rate on the microstructure evolution from the nucleation to a full solidification stage. 2.6 Nucleation Processes The solidification of an alloy is initiated by the nucleation of the first solid within a liquid melt. This process can be captured using thermal analysis techniques by placing a thermocouple into the liquid metal as the metal cools. Figure 2.1 illustrates the cooling curve of a pure metal. The point “A” represents the initial pouring temperature of the metal. Point “B” is assumed to be the temperature when nucleation begins. Undercooling must occur for the first nuclei to form. 12  The magnitude of undercooling is determined by the extent of homogeneous or heterogeneous nucleation. This will be discussed in further detail in section 2.6.2. Between points “B” and “C”, the latent heat evolution (also referred to as recalescence) occurs, while at  point “C” grain growth begins and the grains continue to grow until the material is completely solid at point “D”. The solid material continues to cool to ambient room temperature (point “E”).  Figure 2.1: Representative cooling curve for a pure metal [31] Cooling curves provide helpful information to determine the nucleation temperature for a given pure metal. In alloy systems, the cooling curve looks different. Cooling curves for alloy systems contain a range of temperatures (solidification range) where solid and liquid phases can exist in equilibrium. The thermodynamics of binary alloys allow for two (or more) phases to exist with one of the phases being liquid. As the temperature begins to rise or lower, the volume fraction of solid phases will lower or rise, respectively. The correlation between the cooling curve and phase diagram of a binary alloy is seen in  Figure 2.2 below.                A cooling curve is unique for a given alloy; as such, helpful information can be gathered to understand solidification behaviour for a specific alloy.  Figure 2.2: Cooling curve [left] for a binary phase diagram [right] [70]  13  2.6.1 Homogeneous Nucleation Homogeneous nucleation occurs when a solid spherical particle of a pure element is formed within a liquid of the same element. The size of the solid particle becomes larger as the temperature falls below the melting point of the material [12]. The initiation of homogeneous nucleation involves the formation of ordered regions of the solid phase, which continually form and dissociate [12]. These ordered regions require high orders of undercooling energy to begin forming nuclei. The energies considered during undercooling are the surface energy (interfacial energy) and the bulk free energy (volume free energy). The bulk free energy is always a negative term and is a function of temperature and radius of the nucleus. In contrast, the surface energy is always positive. Figure 2.3 shows the surface energy as the blue curve and the bulk free energy as the red curve. The sum of the two energies is shown with the green curve and is often termed the total energy change. The total energy change introduces a new term, r*, known as the critical radius. Nuclei with a radius smaller than the critical radius will not have enough free energy to survive in the liquid metal and will dissolve. On the other hand, nuclei with larger radii will grow and become stable [7], [12].  It is important to note that when the liquid metal is at lower temperature, a smaller critical radius is can begin nucleation. Conversely, when the liquid metal is at a higher temperature, a larger critical radius will be required to begin nucleation.  Figure 2.3: Free energy change during homogeneous nucleation [32] 2.6.2 Heterogeneous Nucleation Heterogeneous nucleation occurs from foreign surface within a liquid metal. Foreign surfaces include: inclusions, oxides, mold walls, and grain refining particles [12]. Undercooling 14  is still considered on the heterogeneous cooling curve; however, at a much lower magnitude than in the case of homogeneous nucleation. This is illustrated in Figure 2.4 below.   Figure 2.4: Free energy comparison for homogeneous (blue) and heterogeneous (red) nucleation [33] Rather than considering the nucleation of a sphere as in homogeneous nucleation, a spherical cap is used to show the various interfacial energies between atomic planes during heterogeneous nucleation. Figure 2.5 shows these interfacial energies.   Figure 2.5: Surface energies during heterogeneous nucleation [12] Where, γβ-L is the nucleus-liquid interfacial energy, γα-β is the substrate- nucleus interfacial energy, γα-L is the substrate-liquid interfacial energy, and θ is the wetting angle. The relationship between the interface energies is represented in Equation 2.1 below. γα−L = γα−β + γβ−L cos 𝜃 Equation 2.1: Interfacial energies for heterogeneous nucleation [12] The ideal case for heterogeneous nucleation is also referred to as complete wetting and occurs when  is zero. With this geometric advantage, the critical radius can be made with fewer atoms; allowing the nucleus to survive and grow into a grain. It can be concluded that the 15  nucleation process has a large role in determining the final grain structure and effectively the material’s properties. 2.7 Grain Refinement Grain refinement is a process wherein the microstructure of the metal during solidification is manipulated in a way to produce a small and equiaxed grain structure at the end of the solidification process [12].  Grain refinement can be achieved by controlling the cooling rate during solidification using integrated cooling systems and/or by inoculating the liquid metal with stable nucleating particles before pouring the liquid metal into a mold [12], [34]. The latter process introduces particles which offer heterogeneous nucleation sites and effectively reduce the required Gibbs free energy required for grain growth [12], [18]. The addition of inoculating particles into the molten material is relatively cost effective and is employed in industry to improve the microstructure and strength of castings [8], [14]. However, the density difference between the inoculating particles and the molten alloy inherently present challenges to achieve a uniform dispersion of the inoculants. In an effort to combat this issue, manufacturers have developed secondary alloys called “Master Alloys”, which carry the inoculant in a compatible matrix with the alloy to be grain refined.  2.8 Master Alloys The use of master alloys allows manufacturers to fine-tune the properties of their castings. Aluminum-based master alloys typically contain, by weight, 90-99% (or 90-99wt%) Aluminum with the balance being the inoculating particles. This creates near unity of densities between the master alloy and the molten aluminum matrix. Thus, the master alloy can be suspended inside of the molten aluminum, where it can dissolve and release the inoculating particles [20], [35], [36].  The inoculating particles used for changing the grain structures in aluminum metal casting often include Titanium Boride and Titanium Carbide. 2.8.1 Titanium-Boride Grain Refiners Titanium Boride, TiB2, is a ceramic material, which possesses excellent resistance to oxidation, great heat conductivity and a high hardness [37]. The effect of TiB2 on grain refinement has been heavily researched due to its proven empirical nucleating ability. Mohanty et al. introduced synthetic Boron crystals to an Al-Si foundry alloy in an attempt to facilitate the nucleation of the α-Al phase. They discovered that TiB2 was not responsible for nucleation directly; however, it provided a substrate for precipitation of TiAl3, which subsequently 16  nucleated the α-Al phase [19]. Wang et al. performed a similar study, where different Al-Ti-B master alloys were produced and tested in a commercially pure aluminum. They concluded that a layer of TiAl3 dissolved into the aluminum via peritectic reaction to nucleate the α-Al phase [38]. Further researchers observed that Ti-B based master alloys also contain amorphous (All-xTix)B2 particles and form a solid solution. Solute particles such as silicon and chromium, can have an adverse effect on the grain refinement of aluminum alloys. Namely, by altering the chemistry of the inoculating particles, nucleating potency can be retarded, thus making them incapable of supporting the α-Al phase formation [38]. Moreover, agglomeration of TiAl3 and TiB2 reduces the number of active Ti-rich particles in the melt, and thus grain refinement is poisoned. These unfavorable effects prompt further research in stable inoculants such as Titanium Carbide. 2.8.2 Titanium-Carbide Based Grain Refiners The following section discusses the mechanisms reported in the literature on titanium-carbon grain refinement of aluminum alloys. Additionally, the synthesis of titanium carbide and properties as a grain refiner are also discussed. 2.8.2.1 Titanium Titanium has been added to aluminum castings for many years to improve various mechanical, physical and chemical properties. The similar lattice structure (FCC), atomic radii, and electronegativity between aluminum and titanium suggest that a formation of a solid solution should readily take place [39].  Table 2.5 gives a comparison of select parameters for titanium and aluminum [7]. Parameter Titanium Aluminum Lattice structure FCC FCC Atomic Radius (pm) 147 143 Electronegativity (Pauling) 1.54 1.61 Table 2.5: Atomic parameters of titanium and aluminum [7] The Aluminum-Titanium phase diagram in Figure 2.6 shows the effect of small titanium additions on the liquidus temperature of an Al alloy. The increase in the melting temperature from 660.1°C to 665°C results from a peritectic reaction.  17    When Ti is added to liquid Al, the TiAl3 phase forms at the Ti-Al interface [13], [20]. TiAl3 crystals begin to dissolve when put into contact with aluminum, which increases the solidification temperature of the alloy [18], [34], [40]. The creation of a localized phase which has a melting point above the base aluminum or aluminum alloy fosters ideal conditions for nucleation of aluminum grains when the temperature of the solution begins to drop. However, TiAl3 is fully soluble in aluminum, if the alloy is held at high temperature for too long. This fading effectively disables its ability to nucleate grains during long hold times. This issue has been partially resolved with the addition of a carbon source to the master alloy, which has enabled grain refinement of aluminum with titanium [36], [41], [42].   2.8.2.2 Carbon Black Carbon black (CB) is one of the polymorphs of Carbon and contains an amorphous structure that is stable at high temperatures. The structure and characteristics of CB depend on the manufacturing process. Table 2.6 provides a summary of the most common CB manufacturing methods [43].     Figure 2.6: Aluminum-titanium binary phase diagram  [14] 18  Process Details Furnace Black • Partially combusted petroleum products into a furnace.  • Improved control of amorphous particle size and structure.  • Ideal for mass production. Channel • Partially combusted natural gas toward a channel steel  • Ability to create functional groups on the surface for later processing Acetylene Black • Biproduct of decomposition of acetylene gas  • Manufacturing process allows for very high purity  Lampblack • Soot collected from burning of oils or pinewood Table 2.6: Methods for processing and utilizing carbon black [43]  Carbon black is defined and categorized by its particle size and structure. Carbon black has a high surface area to volume ratio, which makes it amenable to react with titanium. Literature suggests that carbon black with large particle sizes typically contains a more organized surface structure with fewer available activation sites [44]. The work of In-Shup et al. [45] demonstrated that a reaction between titanium hydride (TiH2) and carbon black enables the formation of sub-micron TiC particles. The reaction was achieved by heat treating ball milled blends of TiH2 with carbon black in the range of 600-1000°C. Furthermore, in another study, longer milling times lowered the required synthesizing temperature [46]. The production of TiC inoculating particles from Ti and CB could have significant ramifications in the economical production TiC-based master alloys from reduced material and energy costs. 2.8.2.3 Titanium Carbide The synthesis of Titanium carbide powder (US Patent Number: 5,417,952) begins by adding  Titanium dioxide powder with an average particle size of 0.3µm and surface area of 50m2/g into a reactor where a low density and highly porous carbon coating (C3H6) is applied at 420oC. The coating is applied to the Titanium dioxide until the particles are ~34wt% carbon. The carbon-coated titanium dioxide particles are then heated to 1,450 oC for one hour in a continuously flowing argon environment.   Titanium carbide has a melting point of  2,937oC and a face-centered cubic (FCC) lattice structure (ao=0.4329nm) [47]. Improved stability at elevated temperatures, when compared to titanium, has researchers interested in Titanium Carbide’s abilities to act as a nucleation site for 19  grain refinement purposes. Cibula’s 1949 ‘Carbide Theory’ suggests that upon the addition of titanium into aluminum, segregation of residual carbon results in the spontaneous formation of TiC [48]. Although offering little success, this theory has perpetuated the works of numerous researchers to investigate TiC’s ability to nucleate grains in aluminum [19], [34], [41], [42], [49]. Banerji et al. have developed methods to produce TiC via Al-Ti-C master alloys [48], [50], [51]. In one study, the preparation of an Al-Ti-C master alloy was put forward by the addition of graphite and amorphous carbon into an Al-Ti binary alloy with mechanical stirring [51]. The carbon particles created two different reactions. The carbon-titanium reaction resulted in the successful formation of TiC. However, the carbon-aluminum reaction has resulted in the formation of Al4C3, an unfavorable phase responsible for poisoning the TiC and effectively hindering its potency to nucleate grains, as reported in [51]. Wang et al. tested a their Al-5Ti-0.75C master alloy in commercially pure Al. The master alloy was produced by combining pure Ti and Al with Al-5C alloy using a melt reaction method, and contained a homogeneous distribution of α-Al (matrix), TiAl3 (fine particles), and TiC (plate-like) phases throughout the master alloy [36].  No poisoning phases formed and therefore the fabrication method successfully produced an Al-Ti-C master alloy. Leon et al.  [52] performed a sessile drop experiment to find the wettability of TiC with commercial aluminum alloys (1010, 2024, 6061,and 7075). The authors suggested that the interaction between molten Al and TiC substrate formed the unfavorable Al4C3 phase at the Al-TiC interface. This unfavorable Al4C3 phase had adverse effects on wetting. The reactions between molten 2024 alloy and TiC substrate exhibited the reduced presence of Al4C3 that resulted in the increased adhesion  of 2024 alloy to TiC [52]. The improved adhesion was correlated to the observation of CuAl2 intermetallics residing at the 2024/TiC interface. However, the presence of the intermetallic phase did not directly contribute to improved adhesion; rather, it did delay the formation of Al4C3 [52]. This is consistent with the work of Froumin et al. who reported that the presence of Cu decreased the solubility of Ti in molten Al [53]. This suggests that titanium-carbon grain refiners can serve as stable nuclei to heterogeneously nucleate grains in aluminum B319. The work of Kumar and Bichler [34] focused on the addition of titanium carbide powder to aluminum B319. They discovered that the TiC powder did successfully refine the grains and improve the mechanical properties of the as-cast alloy. The improvement in strength was attributed to refined α-Al dendrites and improved segregation of secondary phases [34].  20  2.9 Effect of Grain Refiners on 319 Alloys The following sections review the most up-to-date literature pertinent to historical development of grain refiners for industrial applications and castings with the 319 aluminum alloys. 2.9.1 Alloy Microstructure and Morphology The B319 alloy contains a combination of eutectic phases and intermetallic particles suspended in an aluminum matrix. Figure 2.7 shows the etched (A) macrostructure, (B) microstructure, and (C) unetched SEM image of as-cast aluminum B319.  The alloy’s grains experience dendritic growth with Si-rich phases segregating between dendrite arms; seen as blue in Figure 2.7 (B). Figure 2.7 (C) shows the silicon particles under high magnification and these appear as dark regions. The large white fishbone structures are the AlMnFeSi intermetallic particles. It is generally accepted that inoculating particles can change the size and morphology of the secondary phases in addition to reducing the secondary dendrite arm spacing (SDAS) of the primary phase grains.  Kumar and Bichler successfully reduced SDAS by ~ 20%, when 0.03wt% TiC powder was added to the B319 alloy [34]. Although a 20% reduction in SDAS was achieved, no observations to the morphology of the interdendritic eutectic and secondary phases was noted. Shabestari et al. added a Al-5Ti-1B master alloy to molten 319 alloy [40]. Using differential thermal analysis (DTA), the results suggested that the addition of the master alloy increased the solidification range (range of temperatures from initial solid to final solid) and total solidification time (time from initial solid to final solid)  [40]. The inoculating particles had a considerable effect on the solidification ranges of the α-Al dendrites and the Al-Si eutectic, but had only a A Figure 2.7: As-cast Aluminum B319: (A) etched macrostructure, (B) etched microstructure, (C) SEM of intermetallics Bb C 21  limited effect on the Cu-eutectic phase. Additionally, the particles reduced the solidification time to develop the dendrite network and the Al-Si eutectic formation, but caused no significant changes in the Cu phase. As a result, the particles had a little to no effect on the development of the Cu phases.  2.9.2 Fluidity As discussed in section 2.3.1, silicon is added to the 319 alloy to improve the alloy’s fluidity and resistance to hot tearing. Dahle et al. tested the effect of a commercial Al-5Ti-1B grain refiner on the fluidity of Al-7Si-1Mg and Al-11Si-1Mg alloys using a spiral sand mold [54]. The study concluded that grain refinement postponed dendrite coherency to a higher fraction of solid. Thus, alloys with a higher fraction solid at the dendrite coherency point offered greater fluidity [54].  This was consistent with the findings of Marchwica [26]. Therefore, through grain refinement, a material is able to behave as a liquid for longer durations during solidification. Veldman et al. evaluated the dendrite coherency of various Al-Si-Cu alloys using rheological and thermal analysis methods [55]. They discovered that the coherency fraction solid was governed by the silicon content. It was observed that increasing the Si content lowered the coherency fraction solid. The study concluded that the coherency fraction solid was best correlated to Si’s effect on dendritic morphology and had little relation to the alloy’s grain size or grain growth velocity. This was not consistent with the findings of  Chai et al. [56] who concluded that dendrite coherency is governed by dendrite growth and grain size. They claimed that the dendrite coherency was inversely proportional to the ratio (V/d) where (V) was the dendrite growth velocity and (d) was the grain size diameter. This suggests that if the solidification conditions are such that growth velocity decreases or the grain size increases, the dendrite coherency point would be delayed during solidification and subsequently the fluidity of the metal would increase.   2.9.3 Mechanical Properties The strength of an alloy can be generally related to its grain size. This relationship is represented using the Hall-Petch Equation as follows [17]: 𝜎𝑜 = 𝜎𝑖 +𝑘√𝐷 Equation 2.2: Hall-Petch equation [17] 22  Where, o: is the ultimate tensile strength of the material. i: starting stress for dislocation movement. k: strengthening coefficient. D: average grain size. This relationship is helpful in determining the strength of a material based on its grain size; however, in multicomponent alloys, secondary phases also contribute significantly to the hardness via precipitation hardening [27]. Due to its complex interactions, the effect of grain refiners on the segregation of secondary phases has not been resolved in the literature yet. Studies have reported on the effect of inoculating particles to improve the mechanical properties of  binary alloys like Al-Cu and Al-Si [19], [36], [57]. Wang et al. tested the effect of their Al-5Ti-0.75C master alloy in an Al-5Cu alloy casting [36], [41]. The addition of their master alloy to the Al-5Cu casting reduced the average grain size by 50% and subsequently increased the ultimate tensile strength (UTS) from 206.8Ma to 260.4MPa [36]. The results were correlated to a fine equiaxed grain structure and dispersed distribution of secondary Cu phases [36]. Rao et al. tested their Al-5Ti-2C-15Sr master alloy in a hypoeutectic Al-Si alloy (LM25) casting [57]. The effect of the master alloy on the LM25’s microstructure and wear behavior was examined. The master alloy reduced the SDAS by ~63% and improved the wear resistance from 250m/mm3 to 365m/mm3. The results came from a combination of refinement of α-Al from TiC particles and modification of Si eutectic from Strontium (Sr). Sr additions modified the Si eutectic morphology from large (coarse) to small (fine) needle-like structure.  2.10 Spark Plasma Sintering (SPS) Since the early 1960’s, spark plasma sintering and electric-field assisted sintering processes have seen major technological advancements and industrial adaptation [58]. The ability to sinter a variety of material classes has gained the interest of materials engineers and scientists worldwide. The SPS process presses powdered particles via hydraulic rams inside of an electrically conductive die. Simultaneously, the rams apply joule heating (ohmic resistance heating) due to the DC electric current passing through the powder. Typically, SPS processing is done in an inert environment. The inert atmosphere inhibits the growth of adverse oxide layers forming on the particle surfaces.  The combination of good particle contact and joule heating allows for rapid and relatively homogeneous heating throughout the powdered material to facilitate bonding of neighboring powder particles [59]. Benefits of this technique includes rapid and controlled heating, minimal gas and impurity inclusion at particle-particle interfaces. Producing materials in this way allows 23  researchers to produce materials that are not possible to find in nature. Trzaska et al. conducted a study to understand the mechanisms involved during the densification of TiAl powder during SPS [60].  The study found that no signs of microstructural changes (i.e. no lamellar grains by localized overheating and cooling) were observed during processing. However, high density crystal defects, dislocations, and deformation twins were found at the particle interface while the center of the particles showed very few defects. It was concluded that intense plastic deformation in the contact zones was induced by SPS resulting in bulk twinning deformation. Lazurenko et al. investigated the formation of Al3Ti, Al2Ti, AlTi, and AlTi3 intermetallic phases from sintering Ti and Al powders at 1050oC  [61]. The Al3Ti compound was the first to form a solid-state reaction during the sintering schedule due to its low free energy of formation.    24  Chapter 3: Experimental Procedure This chapter contains a comprehensive overview of the experimental procedures used in this research. The chapter covers the preparation of the master alloys and casting materials. Detailed experimental schedules are included in addition to the methods and instruments used to characterize casting processes and properties of as-cast materials.  3.1 Stages of Experimental Procedure The experiments pertaining to this research were carried out in three stages. Specific details for each stage were as follows:  1) SPS Master Alloy Development: The materials, equipment and procedures used to prepare the SPS’d grain refining master alloys are described. 2) B319 Casting Experiments: Comprehensive overview of casting schedules, equipment and procedures used to fabricate B319 aluminum alloy castings. 3) Material Characterization: Details pertaining to the methods and equipment used to evaluate the properties and microstructure of as-sintered SPS master alloys, as well as as-cast B319 aluminum alloy castings.  3.2 SPS Master Alloy Development All master alloys were prepared at the University of British Columbia’s Spark Plasma Sintering (SPS) laboratory. This section discusses the materials and processing parameters in this task.  3.2.1 Materials 3.2.1.1 Aluminum powder An aluminum (Al) powder (Chemical Abstracts Service (CAS) number 7429-90-5) with 99.5% purity was used as a matrix for all master alloys. The Al powder mesh size was #325 (<40µm). The powder was stored in a vacuum environment to avoid oxidation of the powders prior to sintering.  25  Figure 3.1 A shows a macro image of the powder prior to blending and sintering. The raw powder had a very good fluidity when handling and measuring. Figure 3.1 B shows an SEM image of the individual powder  particles.  3.2.1.2 Titanium powder A titanium (Ti) powder (CAS # 7440-32-6, Alpha Aesar product # 42624), with a 99.5% purity and a mesh size of -325 (<40µm) was also used. Figure 3.2 shows a bulk image of the Ti powder. As can be seen from the images, the Ti powder had a tendency to agglomerate into clusters.  Figure 3.1: Aluminum powder: [A] macro image, [B] SEM micrograph  Figure 3.2: Titanium powder  A B 26  3.2.1.3 Carbon Black powder The Carbon Black (CB) powder (CAS #1333-86-4) was obtained from Alpha Aesar (product # 45527) and was classified as an Acetylene Carbon Black. The powder was 99.9+% pure Carbon and 100% compressed to limit the variation in particle size and structure.  Figure 3.3A below is a sample of the Carbon Black powder used as an additive to the master alloys. The CB powder was free flowing and did not agglomerate. Figure 3.3B shows an SEM image of the particles. 3.2.1.4 Titanium Carbide powder The titanium carbide (TiC) powder (CAS #12070-08-5) was obtained from Alpha Aesar (product #40178) had a 99.5% purity and an average particle size of 2 µm. Figure 3.4 below is a sample of the powder used. The powder was dark gray in color and had a tendency to agglomerate into clusters when transferred between vessels.  Figure 3.3: Acetylene carbon black: [A] macro image, [B] SEM micrograph (50,000x) Figure 3.4: Titanium carbide powder (CAS: 12070-08-5) A B 27  3.2.1.5 Powder Blending and Processing All powder blends were prepared at UBC’s SPS laboratory. Table 3.1 summarizes the various master alloy composites prepared. Al was used in all blends to help promote a uniform distribution of inoculating particles (TiC and TiCB) throughout the B319 alloy during casting. For all experiments where a master alloy was added to treat a B319 aluminum alloy, 0.03wt% of inoculating particles was always added.  Master Alloy Recipe (wt%) Identification Aluminum – Titanium Carbide (SPS) [Low Concentration TiC] 99 Aluminum 1 Titanium Carbide Al-1TiC Aluminum – Titanium Carbide (SPS) [High Concentration TiC] 95 Aluminum 5 Titanium Carbide Al-5TiC Aluminum – Titanium – Carbon Black (SPS) [Low Concentration TiCB] 99 Aluminum 0.9 Titanium 0.1 Carbon Black Al-1TiCB Aluminum – Titanium – Carbon Black (SPS) [High Concentration TiCB] 96 Aluminum 4.6 Titanium 0.4 Carbon Black Al-5TiCB Table 3.1: List of all master alloys developed to refine B319 grain size. In addition to the grain refiners listed in Table 3.1, pure TiC powder was also added to the B319 alloy during the casting experiments. This powder was not subjected to the SPS process.  Further, pure Al powder was sintered via SPS and later used to treat the B319 alloy during casting in a similar way as grain refiners listed in Table 3.1. A summary of the pure metals used are seen in Table 3.2 below. Table 3.2: Control additives  The fabrication of the Al-1TiC, Al-5TiC, Al-1TiCB, and Al-5TiCB master alloys was carried out as follows:  Master Alloy Recipe (wt%) Identification Titanium Carbide (Powder) 100 Titanium Carbide TiC Powder Pure Aluminum (SPS) 100 Aluminum Pure Al 28  Figure 3.6: Powder blend after manual shaking Step 1: A high-precision analytical balance (+/- 0.00005g) (Fisher Scientific Model number: ALF204) was used to weigh ~50g of powders (Figure 3.5).   Step 2: The weighed powders were transferred to a 100mL glass jar and capped with lid and manually shaken to pre-mix the powders (Figure 3.6).     Figure 3.5: Analytical microbalance 29  Figure 3.7: Ball mill canister and zirconia balls Step 3: The powder blends were transferred to a ball mill canister (FRITSCH: 50.9660.00) seen in Figure 3.7 with ~250mL of alcohol (Fisher Scientific: Alcohol Denatured – 85.45% Ethanol, CAS: 64-17-5). Step 4: Six (6) 10 mm diameter Zirconia balls were added into the vials (Figure 3.7).     Step 5: The vial was properly sealed with a lid and clamps and inserted into the FRITSCH: Pulverisette 7 high energy planetary ball mill (Product number 07.5000.00) seen in Figure 3.8. The blending cycle was set at 3 minutes spinning at 150RPM with a 1 minute pause. (Repeated four times).   Figure 3.8: Fritsch ball mill  30  Step 6: The powder blends were transferred to a clean glass petri dish and heated to ~60°C to accelerate alcohol evaporation. This process was completed within ~2 hrs. Once dried, the blends were transferred to a clean jar and sealed. To avoid oxidation of powders, the blends were be sintered as soon as possible after drying. 3.2.1.6 SPS Supplies and Equipment The spark plasma sintering fabrication of master alloys followed powder blending. Prior to sintering, a cylindrical dye and matching punches made from extra fine extruded graphite were cleaned with alcohol and wiped with a lint-free Kimwipe ®.  Two sizes of graphite dies were used to optimize the number of sintering experiments needed to fulfill the required amount of master alloy needed for subsequent casting experiments. The sizes used were: a. ϕ60mm (Used for low concentration master alloys) b. ϕ20mm (Used for high concentration master alloys) The quantity of powder in the die was calculated to yield ~10 mm thick pellet. The ϕ60mm die required approximately 120g of powder blend, while the ϕ20mm die required approximately 10g of powder blend. Figure 3.9: ϕ60mm x 10mm SPS master alloy 31   Figure 3.10: ϕ20mm SPS master alloy (left:Al-5TiC, right Al-5TiCB) 2. The inside of the die was lined with graphoil sheet, as seen in Figure 3.11  3. Appropriately weighed powders were poured into the open end of the die, and covered with graphoil disk and punch. 4. The contained powder assembly was inspected and wrapped with carbon insulting blanket.   Figure 3.11: ϕ20mm  punches, die, and graphoil for SPS 32  The SPS machine used was a Thermal Technologies® 10-3 SPS machine, seen in Figure 3.12 below. Thermal Technologies® 10-3 SPS machine consisted of a sintering chamber, DC power source, hydraulic press and a vacuum system.  Figure 3.12: Thermal technologies 10-3 SPS machine   3.2.1.7 Sintering Schedule The sintering schedule was based on the work of Sweet et al [59]. All sintering schedules were as follows: 1) Die and powder assembly were arranged between the rams with appropriate graphite adaptors ensuring flush contact between all surfaces.  2) A preload of 5MPa was applied to the die assembly. Sintering Chamber DC Power Source Hydraulic Press Vacuum System (not seen in image) 33  3) Chamber was evacuated to 0.02 Pa vacuum to remove any ambient air. Then, the chamber was back-filled with inert Argon gas. 4) Rapid temperature ramp to 150oC and held at 150 ºC for 1 minute 5) Rapid temperature ramp to 300oC and held at 300 ºC for 1 minute 6) Slow temperature ramp to 470oC with matching applied force ramp to 30 MPa and held for 6 minutes. Sintering began at this stage. The slow ramp was required to minimize the chances of the system from overshooting the target temperature. 7) Temperature was ramped down from 470 oC to room temperature in approximately 15 minutes. The thermal and pressure profiles of the sintering schedule are seen in Figure 3.13 below.  Figure 3.13 Thermal and pressure profile of sintering schedule used to produce SPS master alloys Once the die and punch was safe to handle, the punches were removed from the die and the sintered material (SPS pellet) was ejected from the die using a hydraulic press. The graphoil was removed from the pellet and discarded.   34  Figure 3.15: Preheat furnace (euclid CF 510) Figure 3.14: Melting furnace (euclid R-85) 3.3 Casting Experiments All metal casting experiments were performed at UBCO’s metal casting laboratory. Equipment used in the casting lab space consisted of: 1. Melting Furnace (Euclid Model: R-85 in Figure 3.14): Used to superheat aluminum B319 aluminum alloy.   2. Preheating Furnace (Euclid Model CF 510 in Figure 3.15): Used to preheat molds prior to casting.    35  3. Thermocouples (k-Type) a. Small: OMEGA HKMTSS-020E-24 [Sheath Diameter: 0.02”, Tolerance: ±2.2oC] b. Medium: OMEGA KMQSS-062G-24 [Sheath Diameter 0.062”, Tolerance: ±2.2oC] c. Large: OMEGATJ36-CAXL-14G-28[Sheath Diameter 0.25” , Tolerance: ±2.2oC] 4. Silicon Carbide Crucibles a. Small Crucible: Smelko Foundry Products - #3 Bilge b. Large Crucible: Smelko Foundry Products - #60 Bilge 5. Insulating Blanket: Unifrax – Fiberfrax ® Ceramic Fiber Blanket (various sizes cut from stock roll) 6. Heating Tape: OMEGA Ultra-High Temperature Heating Tape (STH051-080) 7. Casting Molds a. TP1: Graphite Store – Medium Extruded Graphite Rod 2”OD (GR060-ROD-2OD) b. Tensile: Medium Extruded Graphite Rod 4”OD (GR060-ROD-4OD) c. Fluidity: Custom mold (Material: 1020 Steel) The B319 alloy was supplied by Nemak Canada in ingot form with each ingot weighing ~13kg. The required amount of material was sectioned from these ingots. 3.3.1 TP1 Mold The TP1 mold is standard means of examining the effect of melt additives on the as-cast alloy microstructure  [13], [34]. The mold design was modified to enable thermal analysis as well.  3.3.1.1 Material The molds were manufactured at UBCO School of Engineering’s machine shop from a ϕ56mm medium extruded graphite rod. The mold cavity was cylindrical with dimensions of ϕ25.4mm x 101.6mm. A small through hole, ϕ3.2mm, was drilled into the bottom of the mold where a sacrificial thermocouple and plug was inserted. The thermocouple was inserted so that its tip was 25±0.5mm from the bottom of the mold cavity (Figure 3.16).  36  Figure 3.17: TP1 mold  assembly wrapped with heating tape and insulted with FiberFrax blanket   The mold was wrapped with a single layer electric resistance heating tape. The mold assembly was then completely wrapped in a single layer of insulting blanket (Figure 3.17).  Figure 3.16: TP1 mold set up [all units in mm] 37  3.3.1.2 TPI Mold Casting: Process Parameters The following process parameters were used for all TP1 mold casting experiments: 1. The heat tape from the mold assembly was plugged into the control unit and the TP1 mold was preheated to 350oC at least 30 minutes before casting. 2. Approximately 150g of B319 alloy was weighed to an accuracy of ±0.05g. The weight was recorded so that the proper amount of master alloy could be added. 3. The appropriate amount of master alloy was sectioned into small pieces using a slow-speed cutting diamond saw.  a. The master alloys were weighed so that 0.03wt% of inoculating particles were added to the entire casting.  b. The SPS pellets were sectioned into small disks to reduce the required time to melt. c. The sectioned master alloy was preheated to 100oC to remove any moisture. 4. The proportioned B319 was placed inside of the small silicon carbide crucible that was put into the melting furnace 5. The B319 alloy was heated as follows: a. 710oC for unrefined and powder experiments. b. 710oC for high concentration master alloy and pure aluminum experiments. c. 730oC for low concentration master alloy experiments. 6. Once the alloy was liquid, the melt surface was skimmed to remove an oxide layer which formed during holding the B319 alloy in the liquid state. Once skimmed, protective CO2 gas was placed on top of the newly exposed liquid B319 to avoid further oxidation. 7. The data acquisition (DAQ) system was used to record the temperature of the center thermocouple of the TP1 mold cavity.  8. The master alloy was added to the melt and stirred for 30 seconds, followed by a 2 minute holding period. a. After mixing of the high concentration and pure aluminum SPS pellets, the melt temperature dropped from 710oC to 700±5 oC. b.  After mixing of the low concentration SPS pellets, the melt temperature dropped from 730oC to 700±5 oC. 9. After the 2 minute holding period, the melt was removed from the furnace and poured into the mold at 710±5 oC. 38  10. The DAQ recorded the temperature of the melt during solidification for 10 minutes. a. Temperature readings were recorded every 0.55 seconds (~1.8Hz). 11. The DAQ data was saved, thermocouple was cut, and the casting was ejected from the mold 10±1 minutes after pouring. 3.3.2 Tensile Mold Casting 3.3.2.1 Material The molds were manufactured in UBCO’s machine shop with ϕ102mm (4”) medium extruded graphite rod. The design used a cylindrical cavity with dimensions ϕ25.4mm x 102mm, leaving a uniform wall thickness of 20mm. The mold had an open riser to feed the mold cavity during solidification.  3.3.2.2 Process Parameters The following process parameters were used for all tensile mold casting experiments: 1. The entire mold was placed in the preheating furnace and heated to 350oC for a minimum of one hour to ensure uniform mold temperature.  2. Approximately 1300g of B319 alloy was sectioned from the ingot and weighed to an accuracy of ±0.05g. The weight was recorded so that the proper amount of master alloy could be added.  a. The master alloys were weighed so that 0.03wt% of inoculating particles were added to the entire casting.  3. The proportioned B319 alloy was placed inside of the large silicon carbide crucible that was inserted into the melting furnace. 4. The B319 alloy was heated with temperature measurements taken using the large thermocouple. a. 710oC for unrefined  experiments. b. 740oC for high concentration master alloy experiments. c. 730oC for low concentration master alloy experiments. 5. Using a ladle, the melt was skimmed to remove the oxide layer. Once skimmed, protective CO2 gas was placed on top of the newly exposed liquid B319 alloy to avoid further oxidation. 6. The master alloy was added to the melt and stirred for 30 seconds followed by a 2 minute holding period. 39  a. After mixing of the high concentration SPS pellets, the melt temperature dropped from 740oC to 720±5 oC. b.  After mixing of the low concentration SPS pellets, the melt temperature dropped from 750oC to 720±5 oC. c. Protective CO2 gas was applied during stirring. 7. The melt was removed from the furnace and skimmed with a ladle. 8. The mold was removed from the oven and molten B319 alloy was poured at 710±5 oC.  9. Casting was ejected from the mold 10±1 minutes after pouring. 3.3.3 Fluidity Mold Casting 3.3.3.1 Fluidity Mold: Material The mold consisted of a cope and drag with integrated down sprue fabricated from a 1020 steel material. The cavity was a spiral with through holes indexed every 102mm a along the spiral’s total running length of 1,118mm. Medium size thermocouples were inserted in the holes along the length of the spiral.  The mold cavity was coated with Jig-A-Loo Graphite Extreme Lubricant (P/N:038-1502-2) to avoid liquid metal erosion of the mold, as well as assist with the ejection of the casting from the mold.  Figure 3.18 below shows a computer-aided drawing (CAD) of the fluidity mold..  Figure 3.18: 3D CAD of fluidity mold 40   A stopper (Figure 3.19) was fabricated to plug the ingate of the down sprue. This allowed the molten aluminum to fill the pouring cup to a designated height for uniform pressure head during experiments. The stopper was fabricated with a ϕ1/2” carriage bolt welded to a rebar handle with a medium thermocouple rigged to the assembly. The thermocouple measured the melt temperature at the moment when the stopper was pulled and the molten metal began to flow.     3.3.3.2 Fluidity Mold: Process Parameters The following process parameter were used for all fluidity mold casting experiments: 1. The entire fluidity mold was placed in the preheating furnace and heated to 315oC for a minimum of one hour to ensure uniform temperature throughout.  2. Approximately 600g of B319 alloy was sectioned from the ingot using a gravity-fed band saw and weighed to an accuracy of ±0.05g. The weight was recorded so that the proper amount of master alloy could be added. Figure 3.19: Stopper and medium thermocouple  41  3. The appropriate amount of master alloy was sectioned into small disks using a slow-speed cutting diamond saw.  a. The sectioned master alloys were weighed so that 0.03wt% of inoculating particles were added to the entire casting.  b. The SPS pellets were sectioned into small disks (~2mm thick) to reduce the required time to melt. c. The sectioned master alloy was preheated to 100oC to remove any moisture. 4. The proportioned B319 alloy was placed inside of the small silicon carbide crucible that was put into the melting furnace. 5. The B319 alloy was heated as follows: a. 760oC for unrefined  experiments. b. 770oC for high concentration master alloy experiments. c. 780oC for low concentration master alloy experiments.  6. Using a ladle, the melt was skimmed to remove an oxide layer which formed during holding the B319 alloy in the liquid state. Once skimmed, protective CO2 gas was placed on top of the newly exposed liquid B319 alloy to avoid further oxidation. 7. The master alloy was added to the liquid B319 alloy and stirred with the large thermocouple for 30 seconds followed by a 2 minute holding period. a. After addition of the high concentration SPS pellets, the melt temperature dropped from 770oC to 735±5 oC. b.  After addition of the low concentration SPS pellets, the melt temperature dropped from 780oC to 735±5 oC. c. Protected CO2 gas was applied during stirring. 8. 30 seconds after the stirring, the DAQ was triggered to record the temperature along the length of the spiral in addition to the reservoir temperature. At the same time, the mold assembly was removed from the furnace and wrapped with fiberfrax insulating blanket.  9. 2 minutes after stirring, the melt was removed from the furnace and set up to pour into the reservoir of the fluidity mold. a. During the set up, the mold temperature had dropped from 315±5oC to 300±5oC. 10. The stopper was inserted into the down sprue to plug the ingate while the melt was poured at 740±5 oC. The down sprue was filled to the top in under 2 seconds and the stopper was pulled allowing the B319 alloy to flow. 42  11. Casting was ejected from the mold 10±1 minutes after pouring. 3.4 Material Characterization  This section describes in detail the various methods used to characterize the as-cast material.  3.4.1 Microstructure Analysis The fabricated castings were examined with Optical Microscopy (OM), Scanning Electron Microscopy and Energy Dispersive Spectroscopy techniques.  3.4.1.1 Sample Preparation for Microscopy Representative bulk casting samples were carefully extracted with a band saw or lathe. Careful attention was paid to the sample’s temperature to ensure no localized overheating occurred. Smaller samples were sectioned using a slow-speed cutting diamond saw. A two-part fast cure acrylic supplied by MetLab was mixed in accordance with manufacturer’s instructions and the samples for microstructure analysis were embedded in the acrylic resin. After the acrylic resin cured, the samples were polished to a mirror finish with metallographic polishing consumables using the schedule found in Table 3.3.  Step MetLab Polishing Consumables Approximate Time (sec) 1 200 grit 45  2 400 grit 20 3 600 grit 20 4 800 grit 30 5 9 µm diamond paste and diamond suspension with designated 9 µm polishing pad. 45 6 3 µm diamond paste and diamond suspension with designated 3 µm polishing pad. 90 7 1 µm diamond paste and diamond suspension with designated 1 µm polishing pad. 120 Table 3.3: Polishing schedule for B319 alloy samples 3.4.1.2 Optical Characterization Optical imaging was carried out using three methods:  • DSLR Camera (Canon Rebel T1i) with Macro Lens (Tamron 90mm Macro 55). • Ancansco Stereo Microscope (1x-4x). • Ziess AxioVert Microscope (5x, 10x, 50x, 100x). 43  The equipment was connected to a computer where the images were analyzed with a Buehler Omnimet and/or ImageJ software. Each software generated digital images that were used for measurement of micro- and macroscopic features of the as-cast microstructure. A minimum of 50 grain size measurements per sample were taken to develop statistical results.  3.4.1.2.1 Etching Etching of sample surfaces was performed under the supervision of Mrs. Michelle Tofteland in UBCO’s chemistry laboratory in order to reveal the material’s grain size. Table 3.4 below lists pertinent details of the etchant used in this study.  Table 3.4: Etchants used to reveal as-cast microstructure 3.4.1.3 Scanning Electron Microscopy (SEM) and Energy-Dispersive Spectroscopy (EDS) All sample preparation for electron microscopy, including sputter coating, sample grounding, and stage alignment were done under the supervision of Mr. David Arkinstall at UBCO’s SEM lab. A Tescan ® Mira3 XMU field emission Scanning Electron Microscope (10x – 500,000x magnification) (Figure 3.20) was used for all electron beam imaging. An Oxford ® Aztec X-MAX EDS attachment and supporting Oxford® Aztec software were used for chemical data processing.   Figure 3.20: SEM/EDS lab station Etchant Chemical Sample Immersion Time (sec) Purpose Keller’s Reagent • 95ml H2O (Distilled Water) • 2.5ml HNO3 (Nitric Acid) • 1.5ml HCl (Hydrochloric Acid) • 1.0ml HF (Hydrofluoric Acid)  20 Evaluate Grain Size 44  Figure 3.21: Spiral fluidity casting on polar graph paper 3.4.2 Fluidity Analysis A relative measure of the grain refiner’s effect on the B319 alloy’s fluidity was done by measuring the flow length of the as-cast spiral. Each casting was photographed with a polar graph paper background and archived for later evaluation – this can be seen in Figure 3.21.     Figure 3.21 also shows the datum “A” where the spiral lengths were measured from. The thermocouples indexed along the spiral’s length were used for thermal analysis. Each spiral was sectioned ~13mm from the spiral end tip and 102mm from the datum where the first thermocouple was located. These sectioned samples were polished and evaluated with optical and SEM/EDS microscopy.  3.4.3 Thermal Analysis The data collected from thermal analysis was plotted on a Temperature (oC) vs. Time (s) graph. The first derivative curve [i.e., solidification rate (oC/s) vs. Time (s)] was also. When used together, these two graphs enabled the determination of the following solidification characteristics (Table 3.5):   A Spiral Tip First (120mm)  Thermocouple Site 45  Symbol [units] Description Detail t (N,α) [s] α-Al nucleation undercooling time Measured time between T(N, α) and when rate was ~0 (oC/s) T(N, α) [C] α-Al nucleation undercooling temperature Measured at the first climb of the rate curve T(Min, α) [C] Minimum α-Al undercooling temperature  The lowest temperature observed during α-Al nucleation T(G, α) [C] α-Al dendrite growth temperature Measured when rate was ~0 (oC/s).  CR(α) [C/s] Cooling rate before first α-Al nucleation Measured between t=0s and first α-Al nucleation T(N,Si) [C] Si eutectic nucleation temperature Measured at the second climb of the rate curve T(N,Cu) [C] Cu eutectic nucleation temperature Measured at the third climb of the rate curve CR(Total) [C/s] Total cooling rate from initial to final solid = 𝑇(𝑁,𝛼) − 𝑇(𝑆) t(S) T(S) [C] Solidus temperature Assumed to be 499oC t(S) [s] Time at solidus Time when casting reached 499oC Solid Time [s] Time between initial and final solid = t (N,α) - t (S) Freezing Range [C] Temperature between initial and final solid = T(N,α) - T (S) Table 3.5: Solidification characteristics of B319        Figure 3.22 below shows the characteristics listed in Table 3.5 from a baseline casting experiment using unrefined B319. 46  Figure 3.23:Tensile sample extraction location  Figure 3.22: Solidification characteristics represented on a cooling curve 3.4.4 Mechanical Properties The following section provides an outline of tensile and hardness testing performed on the as-cast materials.  3.4.4.1 Tensile Testing All tensile tests were performed at UBCO’s high head lab under the supervision of Mr. Alec Smith. Using the graphite mold (discussed in section 3.3.2.1), tensile samples were extracted from the outer-most region of the mold’s cross section as seen in Figure 3.23.    All tensile samples were fabricated in UBCO’s machine shop using a Manual Lathe for turning cylindrical samples followed by the use of the HASS CNC Lathe to generate the ‘Dog-Bone’ profile. A light depth of cut (<0.03”) was used to minimize stress and heating of the -40-35-30-25-20-15-10-5054505005506006507000 5 10 15 20 25 30 35 40Rate (C/s)Temp (C)Time (s)Unrefined  B319Cooling RateTemperaturet&T (N,α) [s & C] t & T (Min, α) [s & C] CR (α)  t & T (G, α) [s & C] t & T (N, Cu) [s & C] t & T (N, Si) [s & C] t & T (S) [s & C] CR (Total)  Solid Time [s] Freezing Range 47  Figure 3.24: Tensile sample dimensions (dimensions in “mm”) Figure 3.25: Instron tensile machine sample. The sample profile was cut in accordance with ASTM B557M-15 “Standard Test Methods for Tension Testing wrought and Cast Aluminum- and Magnesium- Alloy Products” [62]. Figure 3.24 shows the dimensions and tolerances of the tensile samples.      All tensile tests were done using an Instron B557M-15 (250kN) machine seen in Figure 3.25.   48    The test parameters (Table 3.6) were selected in accordance with the ASTM B557M-15 standard. Table 3.6: Parameters used for tensile testing After testing, if inclusions and/or porosity were found at the fracture surface, a retest was performed.  3.4.4.2 Hardness Hardness testing was completed at UBCO in the high-head laboratory using a Rockwell hardness testing machine. All testing was done in accordance with the ASTM E18-15 Rockwell C Scale [HRC] (Diamond indenter and 150kgf Test Force).  Figure 3.26 shows the hardness tester used.  Parameter Value Preload 25N Strain Rate (Cross-Head Velocity) 0.2mm/min Yield Offset  0.2% Failure Criteria 80% Reduction in Load (kN) Figure 3.26: Rockwell hardness testing machine 49  Before the TP1 samples were polished to a mirror finish, the surface was sanded to 400 grit and hardness measurements were collected. Five hardness measurements were taken and averaged on the cross section of each TP1 sample. The units were later converted to the HB scale to enable comparison with published literature.   50  Chapter 4: Results and Discussion In this research, a master alloy was developed via SPS. The master alloy was fully characterized, as well as its effectiveness on grain refining of the aluminum B319 alloy was studied. This chapter provides an in-depth analysis of the obtained results.  4.1 Master Alloy Characterization 4.1.1 Densification of Master Alloy The density of the SPS fabricated master alloys was measured using the Archimedes method, with the results summarized in Table 4.1. The theoretical density values were calculated using the rule of mixtures approach. A sample calculation is provided in Appendix A. Master Alloy Theoretical Density [g/cm3] Actual Density [g/cm3] Densification (Actual / Theoretical) [%] Pure Al 2.700 2.646 96 Al-1TiC 2.722 2.640 – 2.668 97 - 98 Al-5TiC 2.812 2.728 – 2.756 97 - 98 Al-1TiCB 2.715 2.606 – 2.661 96 - 98 Al-5TiCB 2.772 2.661 – 2.717 96 - 98 Table 4.1: Density of SPS’d master alloys  The results in Table 4.1 suggest that the master alloys contained ~2-4% porosity. These porosity and density values are consistent with the work of Sweet et al. [59] who was able to achieve 90-99.5% densification for various sintering experiments using an aluminum powder.  In the present work, achieving a full density was not the main objective of the SPS fabrication of the grain refiners. Since some level of porosity was present in the SPSed master alloy, the increased surface area of grains exposed to a heat flux (originating from the liquid B319 aluminum alloy when the grain refiner was immersed in the aluminum alloy) likely enabled a faster melting of the matrix and a concomitant release of the inoculating particles.  4.1.2 Microstructure of Master Alloy Figure 4.1 shows a representative microstructure of the low concentration Titanium Carbide master alloy (Al-1TiC). The Titanium Carbide (TiC) particles were dispersed throughout the SPSed pellet of the master alloy. The size of the TiC particles varied from 0.5-4 µm, indicating no change in their size as a result of SPS processing, since the raw TiC powder had particles in the same size range. 51  As can be observed in Figure 4.2, the TiC had a tendency to segregate between the much larger Al particles. This was further accentuated as the concentration of TiC in the Al matrix increased. The raw Al powder had a mesh size of 325 (<40µm), while the TiC had an average particle size of 2 µm. Therefore, the physical segregation of the fine TiC on the larger Al particles was the result of the significant particle size difference. In addition to the particle size difference, the relatively hard TiC particles effectively coated the aluminum particles during mixing and were partially embedded into the softer aluminum powder particles, which would have further secured the TiC particles on the Al powder’s grain boundaries.  TiC Al Figure 4.1: SEM image of Al-1TiC master alloy 52  Figure 4.2: TiC particle distribution in Al-1TiC (A) and Al-5TiC (B)   A B 53  In the case of the Titanium and Carbon Black master alloys shown in Figure 4.3, the sintered master alloys contained larger Ti particles with average size of 20µm. Relative to the TiC master alloys, the particle dispersion has improved in these master alloys. Due to the Ti and CB’s size and volume relative to the Al particles, it is evident that the Ti and CB did not coat the Al particles to the same extent as the TiC powder did. Thus, these elements were more homogeneously distributed in the as-sintered master alloy.  Figure 4.3:TiC particle distribution in Al-1TiCB (A) and Al-5TiCB (B) B A 54  4.2 B319 Aluminum Alloy Microstructure Characterization The as-cast B319 aluminum alloy microstructure was evaluated from castings made in the TP-1 and fluidity molds.  The results of grain size measurement and secondary phase analysis are presented in the following sections.  4.2.1 General Microstructure Figure 4.4 reveals that the B319 alloy contained an aluminum matrix solid solution, with eutectic phases and intermetallic particles forming between the primary Al grains.   Figure 4.5 shows the secondary phases and their chemical composition. At the eutectic temperature (577oC) of Al-Si binary alloy, Si has a solubility limit of 12.6wt% and 1.65wt% in Al for liquid and solid phases, respectively [7]. Since Si is not highly soluble in Al, the Si-rich phases precipitating in the B319 alloy were nearly pure silicon [24]. Some Si reacted with Fe and Mn to form complex intermetallic compounds [7]. The distribution and the size of the Si-eutectic particles determines their potency for dispersion strengthening and abrasion resistance [24].  Figure 4.4:Microstructure of unrefined B319 alloy 55       Figure 4.5: SEM/EDS image of secondary phases in B319 microstructure (A) SEM image; (B) Al; (C) Fe; (D) Mn; (E) Si  The dominant secondary phase in the B319 aluminum alloy was the Al-Mn-Fe intermetallic phase, with a large fishbone (or often called “Chinese script”) morphology. The morphology of this phase is directly related to the grain size of the B319 alloy’s matrix, since the buildup of Fe and Mn solutes at the liquid-solid interface was shown to reduce the growth velocity of the primary α-Al grains [63], which effectively slows down grain growth and enhances the time for larger volumes of liquid regions to reside between α-Al grains at the end of solidification [12]. Therefore, in the present work, increasing the number of nucleating α-Al solids during A E B D C 56  solidification due to grain refinement was expected to modify the size and morphology of the Al-Mn-Fe intermetallics.  In addition to the Al-Mn-Fe phase, the B319 alloy also contained a copper-rich intermetallic, with a stoichiometry of CuAl2. This block-like phase, seen in Figure 4.6, was observed throughout the alloy’s microstructure and its morphology was consistent with that reported by other researchers [22], [34].                Spectrum 12 Data Element Wt% Cu 44.89 Al 42.11 O 5.01 C 4.10 Ni 2.25 Si 1.15 Zn 0.49 Total 100 Figure 4.6: EDS chemical analysis of the block-like CuAl2 57  Figure 4.7: Grain structure of unrefined B319 4.2.2  B319 Aluminum Alloy Grain Size Figure 4.7 shows the general grain structure of the as-cast B319 aluminum alloy cast into the TP1 mold. The sample was etched with Keller’ reagent to reveal the grain structure.     The grain structure of the unrefined as-cast B319 aluminum alloy had an equiaxed structure. The B319 aluminum alloy naturally exhibits this type of (equiaxed) grain structure due to two factors. First, nucleation in the B319 alloy begins from solute elements (e.g., silicon), which provide constitutional undercooling and thus restricts grain growth from becoming columnar [12], [34]. Similarly, solutes such as copper, magnesium, iron, and manganese also provide grain growth restriction, due to their enrichment in the solidifying liquid [15], [34].  In addition to the chemical composition, the cooling rate of the alloy also governs the grain size of an as-cast B319 aluminum alloy. A faster cooling rate yields smaller grains and vice versa. As will be discussed in section 4.3.1, the total cooling rate for the unrefined B319 was 4.8 oC/s,  resulting in an average grain size of ~2,500 µm.  Figure 4.8 shows a histogram of the grain size distribution in the unrefined alloy. This distribution had a standard deviation of 662 µm with most measurements found between 1,650 µm and 2,639 µm. 58  Figure 4.8: Grain size histogram for unrefined B319 alloy   0%10%20%30%40%50%60%70%80%90%100%024681012140 330 660 990 1320 1650 1979 2309 2639 2969FrequencyGrain Size (um)Unrefined  B319 Grain Size MeasurementFrequencyCumulative %59  4.2.3 Thermal Analysis The unrefined B319 aluminum alloy was poured into a TP1 graphite mold, which was heated to 350 oC and insulted to maintain the cooling rate between 4.0-7.2 oC/s. The data in Figure 4.9 shows the cooling curve and the first derivative curve for the unrefined B319 aluminum alloy.  Figure 4.9:Cooling curve and first derivative curve for unrefined B319 aluminum alloy Table 4.2 summarizes the solidification temperatures of interest (as defined in section 3.4.3), as obtained from Figure 4.9. Casting t (N,α) [s] T(N, α) [C] T(Min, α) [C] T(G, α) [C] CR(α) [C/s] T(N,Si) [C] T(N,Cu) [C] CR(total) [C/s] T(S) [C] t(S) [s] Solid Time [s] Freezing Range [C] Unrefined 0.37 654.4 607.3 604.4 9.74 583.6 542.6 4.7 499 33.6 33.2 155.4 Table 4.2: Solidification parameters of unrefined B319 aluminum alloy. Limit of error ±2.2oC The cooling curve in Figure 4.9 clearly shows thermal arrests, which corresponded to the α-Al and eutectic phase transformations during solidification. The first derivative (rate) curve indicates three individual climbs that represent the latent heat releases during various phase transformation. The α-Al phases’ nucleation temperature at 654.4oC, followed by Si eutectic at 583.6oC, and Cu eutectic at 542.6oC, respectively. The latent heat releases likely affected the cooling rate of the mushy zone as the material solidified.  -40-35-30-25-20-15-10-5054505005506006507000 5 10 15 20 25 30Rate (C/s)Temp (C)Time (s)Temperature Cooling RateLimits of error ±2.2oC 60  The representative fraction of solid curve in Figure 4.10 was made using the data from thermal analysis of the unrefined B319 aluminum alloy.   Figure 4.10: Fraction of solid curve for unrefined B319 aluminum alloy The fraction solid data of key solidification events is summarized in Table 4.3 below.   α-Al Nucleation (TN,α) α-Al Growth (TG,α) Si-Eutectic (TN,Si) Cu-Eutectic (TN,Cu) Temperature (oC) 654.4 607.3 583. 542.6 Fraction of Solid (%) 0 67 83 98 Table 4.3: Fraction of solid results for unrefined B319 aluminum alloy. Limit of error ±2.2oC According to the table above, a significant portion of solidification occurred during the α-Al phase formation. The final microstructure of the unrefined B319 aluminum alloy discussed in section 4.2.1 showed a well-developed α-Al dendrites with various eutectic and intermetallic phases residing in the interdendritic spaces. This suggests that at the growth temperature the semi-solid solution contained α-Al solids with the remaining 33% of the material being the liquid phase.  4.3 Master Alloy Experiments The following sections first discuss the results of thermal analysis measurements, followed by discussion of microstructure and grain size measurements. This sequence was necessary in order to provide a basis for a subsequent discussion of the effect of master alloys on the B319 00.10.20.30.40.50.60.70.80.91499519539559579599619639659Fraction of SolidTemperature (C)Limits of error ±2.2oC 61  alloy’s grain structure. All thermal analysis results are based on the cooling curves from TP1 mold casting experiments.  4.3.1 Thermal Analysis Figure 4.11 shows the thermal history during the initial stage of solidification of the unrefined and refined alloys.    Figure 4.11: Cooling curve and first derivative; [A] B319 with Al-5TiC and Al-5TiCB addition and; [B]  B319 with Al-1TiC and     Al-1TiCB addition  All of the experiments with the master alloy additions exhibited an increase in the α-Al nucleation and  growth temperatures. When the Al-5TiC and Al-5TiCB was added, the α-Al nucleation temperature increased from 654.4°C  to 659.3°C  and 659.9°C , respectively. The -25-20-15-10-506006056106156206256300 1 2 3 4 5 6 7Rate (C/s)Temperature (C)Times (s)Unrefined TemperatureAl-5TiC TemperatureAl-5TICB TemperatureUnrefined Cooling RateAl-5TiC Cooling RateAl-5TICB Cooling Rate-25-20-15-10-506006056106156206256300 1 2 3 4 5 6 7Time (s)Rate (C/s)Temperature (C)Unrefined TemperatureAl-1TiC TemperatureAl-1TiCB TemperatureUnrefined Cooling RateAl-1TiC Cooling RateAl-1TiCB Cooling RateA B Limits of error ±2.2oC Limits of error ±2.2oC 62  same trend resulted from the addition of Al-1TiC and Al-1TiCB, where the α-Al nucleation temperature increased from 654.4 °C to 660.5°C  and 659.1°C , respectively. The increased Ti concentration, originating from the master alloy addition, in the B319 aluminum alloy was expected to increase the α-Al nucleation temperature because of the peritectic reaction between Al and Ti [14], as was discussed in Figure 2.6 in section 2.7. The high concentration (Al-5TiC and Al-5TiCB) and the low concentration (Al-1TiC and Al-1TiCB) additions were expected to yield comparable α-Al nucleating temperatures because the same inoculating particle wt% was added in all experiments. The difference between the high concentration (Al-5TiC and Al-5TiCB) and the low concentration (Al-1TiC and Al-1TiCB) additions was the additional aluminum. The added aluminum proved to have little influence on the B319 aluminum alloy’s α-Al nucleating temperatures.  By increasing the available nucleation sites with the increased Ti concentration, the required Gibbs Free-Energy needed to nucleate grains was reduced. This effect is best illustrated by comparing the undercooling of the unrefined and the refined B319 aluminum alloys. If the number of nucleation sites has increased, a higher number of grains can nucleate with less required undercooling – resulting in a finer grain structure. In all experiments with SPS master alloy additions, the undercooling was nearly eliminated in comparison to the unrefined B319 aluminum alloy. Further, the resulting refined grain structure from master alloy addition (discussed in section 4.3.3)  was consistent with the reduced undercooling observed in Figure 4.11. Therefore, these observations on undercooling at the nucleation temperature provide evidence that the Ti and TiC were able to facilitate nucleation.  The reduction of undercooling has an influence on the growth of the α-Al phase. The α-Al growth temperature is the steady growth due to release of latent heat of the primary α-Al dendrites.   When the Al-5TiC and Al-5TiCB were added, the α-Al growth temperature increased from 604.4°C to 611.7°C and 610.9°C, respectively. The same trend resulted from the addition of Al-1TiC and Al-1TiCB, where the growth temperature increased from 604.4°C to 612.5°C and 611.3°C, respectively. Shabertari et al. [40] reported increased growth temperature of 319 aluminum alloy after the addition of their Al-5Ti-1B grain refining master alloy. The increased α-Al growth temperature was expected due to the peritectic reaction between Al and Ti. Ti is highly soluble in Al when compared to C [64], and thus it would be expected that addition of the Al-1TiCB and Al-5TiCB master alloys would increase the B319 aluminum’s α-Al growth 63  temperature more than when the Al-1TiC and Al-5TiC were added, since the Ti in the TiC was bound to the carbon. However, the B319 aluminum alloy’s α-Al growth temperature after the addition of TiC was comparable to that of a grain refiner with Ti with CB. Therefore, the CB added to the B319 aluminum alloy may have retarded the reaction between Ti and Al.  Table 4.4 below summarizes the key solidification parameters found in this research.   The cooling rate of the aluminum alloys will influence the solidification parameters listed in Table 4.4 above. Shabestari et al. [65]  performed a study that tested the effect of different cooling rates on the solidification parameters found on the thermal profile of 319 Aluminum alloy. It was found that faster cooling rates may significantly increase the phase transformation temperatures.  Fast cooling rates increase the rate of heat extraction from the melt and allow existing nuclei in the melt to become more activated [65]. Therefore, it is expected that the fast cooling rates used in this study exhibited higher than in published literature [40], [65] where slower cooling rates were employed. Both the Al-5TiC and Al-5TiCB increased the overall cooling rate by increasing the solidification range and reducing the solidification time. The reduction of the solidification time was attributed to the number of growing solids at a given time during solidification. Due to the higher nucleation rate in the grain refined castings, the amount of solid and associated latent heat release increased, thereby achieving complete solidification sooner. Further, literature [14] suggests that the number of growing α-Al grains increases with the reduction of inoculating particle size. This is consistent when comparing the solidification Table 4.4: Summarized solidification parameters of B319 alloy with and without master alloy addition. Limit of error ±2.2oC Casting T(N,alpha) [C] T(Min,alpha) [C] T(G,alpha) [C] T(N,Si) [C] T(N,Cu) [C] CR(total) [C/s] T(S) [C] Solid Time [s] Freezing Range [C] Unrefined 654.4 607.4 604.4 583.6 542.6 4.7 499.0 33.2 155.4 Pure Al 642.7 609.4 611.3 587.6 549.4 4.0 499.0 36.2 143.7 Powder TiC 662.4 607.9 608.5 590.1 549.9 4.1 499.0 39.5 163.4 Al-1TiC 660.5 612.5 611.9 586.0 549.6 5.9 499.0 27.4 161.5 Al-5TiC 659.3 611.7 611.7 588.7 549.3 6.3 499.0 25.3 160.3 Al-1TiCB 659.1 611.3 610.8 588.1 547.7 5.5 499.0 28.9 160.1 Al-5TiCB 659.9 610.9 610.9 587.4 549.1 5.9 499.0 27.1 160.9 64  time of the Al-5TiC to the Al-5TiCB. The Al-5TiC containing smaller inoculating particles had a solidification time of 22.3 seconds, while the Al-5TiCB containing larger particles had a solidification time of 27.1 seconds.  Other works [40], [66] report that inoculating particles used as grain refiners in aluminum alloy castings have the biggest effect on the liquidus phase transformation and associated temperature region, while inoculants have a little effect on the Si and Cu eutectics. However, in the present study the silicon and copper eutectic nucleation temperatures of the B319 aluminum alloy experienced a change with the addition of the master alloys. The silicon eutectic nucleation temperature on average increased 5 °C after the addition of all master alloys. The copper eutectic nucleation temperature on average increased 7 °C after the addition of all master alloys. The same trend was observed by Shabestari et al.. who reported increases of 319 aluminum’s silicon and copper eutectic nucleation temperatures after addition of their Al-5Ti-1B master alloy. In their study, 319 aluminum’s and increased proportional to the level of grain refiner added.  The fraction of solid curves in Figure 4.12 indicate that the addition of the high concentration master alloy had little to no effect on the fraction of solid development in the B319 aluminum alloy.   Figure 4.12: Fraction of solid of B319 with SPS master alloy addition 00.10.20.30.40.50.60.70.80.91499519539559579599619639659Fraction of Solid  f,sTemperature (C)Al-1TiCB Addition Al-1TiC AdditionAl-5TiCB Addition Al-5TiC AdditionLimits of error ±2.2oC 65  The fraction of solid results are inconsistent with the works of Shabestari et al. [40] who reported that the addition of Al-5Ti-1B grain refiners translated the fs curve upwards. Translating the plot upwards was the result of increased nucleation sites allowing for increased enthalpy releases during eutectic reaction [34].  4.3.2 Microstructure Figure 4.13 shows the microstructures of the B319 aluminum alloy after the Al-1TiC and Al-5TiC master alloy additions. In comparison to the microstructure of the unrefined B319 alloy (Figure 4.5), the addition of the Al-1TiC and Al-5TiC master alloys yielded a homogeneous distribution of the intermetallic and eutectic particles. Also, the Si phase between the α-Al matrix reduced in size; however, not to the same extent as the size of the Al-Mn-Fe intermetallics.  Figure 4.13: SEM image of B319 with Al-1TiC [A] and  Al-5TiC [B] (X500) The size of the Al-Mn-Fe intermetallics was greater in the B319 aluminum alloy refined with the Al-1TiC master alloy than in that refined with the Al-5TiC master alloy. As discussed in section 4.3.1, the cooling rate of the B319 aluminum alloy after the addition of Al-5TiC (6.3 oC/s) was slightly higher than that of the B319 aluminum alloy with Al-1TiC (5.9 oC/s), likely contributing to the coarsening of the Al-Mn-Fe intermetallics. The coarsening of these Al-Mn-Fe intermetallics has been correlated to cooling rates in other studies [34].  B A 66  Figure 4.14 shows a representative large Al-Mn-Fe intermetallic that formed after the addition of a pure Al SPS pellet to the B319 aluminum alloy. The amount of Al added in the pellet corresponded to the addition of 3.0wt% of pure Al to the B319 aluminum alloys. The general microstructure of the intermetallics was comparable to that of the intermetallics observed in the unrefined B319 aluminum alloy. Further, the Si eutectic and the CuAl2 phase also remained comparable in size and morphology, when compared to the unrefined B319 aluminum alloy. The thermal analysis data in section 4.3.1 indicates that the addition of pure Al increased the solidification time of the B319 aluminum alloy from 33.2s to 36.2s and reduced the freezing range from 155.4 to 143.7  oC. Consequently, the cooling rate of the B319 alloy decreased from 4.7 to 4.0 oC/s as a result of the pure Al SPS addition.   Figure 4.14: SEM image of B319 with pure Al addition  Figure 4.15 shows the microstructure of the B319 aluminum alloy after the addition of the Al-1TiCB and Al-5TiCB master alloys. The B319 aluminum alloy with the Al-1TiCB and Al-5TiCB had comparable Al-Mn-Fe intermetallic particles. This was expected, since the cooling rate of the alloy system increased only marginally from 4.7 oC/s to 5.5 oC/s after Al-1TiCB addition and to 5.9 oC/s after Al-5TiCB addition.   67   Figure 4.15: SEM image of B319 with Al-1TiCB [A] and Al-5TiCB [B]   4.3.3 Grain Size  Figure 4.16 shows the effect of the master alloys on the general grain structure of the B319 aluminum alloy cast into the TP1 mold. In all of the experiments, 0.03wt% of inoculating particles was added to the entire casting to enable comparison of the present results with the work of Kumar and Bichler [34].    B A 68   Figure 4.17 presents the average grain size of the various experiments. All SPS master alloys (Al-1TiC, Al5TiC, Al-1TiCB, Al5TiCB) reduced the grain size to approximately 40% of the size of the unrefined B319 aluminum alloy.        Unrefined  Pure Al Addition  Powder TiC Addition  Al-1TiCB Addition  Al-1TiC Addition  Al-5TiCB Addition  Al-5TiC Addition Figure 4.16: Comparing grain size between all TP1 castings. (A) Unrefined; (B)Pure Al; (C) Powder TiC: (D) Al-1TiCB; (E) Al-1TiC; (F) Al-5TiCB; (G) Al-5TiC A B C D E F G 69  Figure 4.17: Average grain size measurements of all casting experiments   The variation in the grain size between the castings made with a low and high concentration of TiC and the TiCB master alloys was not statistically significant. The pure aluminum SPS pellet served as a control additive in this research, since the same procedure (i.e., sintering, mixing and casting pouring) was followed for this control additive as well. With the pure Al-SPS pellet, the Al content increased by 3.0 wt%Al, which was comparable to the Al increase when the Al-1TiC and Al-1TiCB grain refiners were added to B319 aluminum alloy. As Figures 4.16 and 4.17 suggest, the addition of the pure Al resulted in a negligible change in the final grain size. This result was expected, since no inoculating particles were present in the master alloy. Since the addition of pure Al to the B319 slightly reduced the cooling rate from 4.7 oC/s to 4.0 oC/s, one would expect larger grains in the final microstructure. However, the average grain size between the unrefined B319 alloy and B319 alloy with pure Al addition was comparable. Therefore, the addition of 3.0wt% Al to the B319 alloy had no statistically significant effect on the final grain size.  70  Similarly, there was no significant difference in the grain size when pure TiC powder was introduced into the B319 aluminum alloy. As often reported in the literature, a powdered TiC would be difficult to effectively distribute when introduced to the molten B319. Further, challenges related to the casting process could have also affected the potency of the pure TiC powder particles in grain refining the B319 aluminum alloy. Although the bulk oxide crust (which formed on the liquid metal during melting of the B319 alloy) was removed prior to pouring the metal into the mold, a thin layer of oxide would readily form on the exposed surface of the liquid B319 aluminum alloy. The powdered TiC had to penetrate through this oxide layer before entering the molten material. This in turn would change the wetting behavior of the TiC particles by the aluminum matrix in the B319 alloy. As reported by Leon et al. [52], an oxide film on inoculating particles prevents the development of a true solid-liquid interface needed for nucleation of potent nucleating particles. According to a study by Wang et al., the distribution, size, and morphology of the inoculating particles influenced the efficiency and stability of traditionally prepared master alloys [41].  Therefore, the similar average grain size measurements from the TiC and TiCB additions in the present work introduce an interesting discussion.  It is known that a carbide may improve the stability of Titanium at elevated temperatures; however, it is unlikely that the addition of Carbon Black enabled the formation of a carbide on the Titanium particles due to relatively low processing times and temperatures.     71  4.3.4 Inoculating Particles Due to the size of the TiC inoculating particles, it was very challenging to find them in the as-cast material. Ti particles were found in the fluidity castings when Al-5TiCB was added to B319 aluminum alloy.  Figure 4.18 shows a cluster of Ti particles within the B319 aluminum alloy. Surrounding the particle is a new phase that was not present in the unrefined B319 casting experiments.   Figure 4.18: Ti particles in the fluidity mold castings  Figure 4.19 shows a closer view of the vicinity of an Ti particle. The linescan data confirmed the identity of the Ti particle in addition to elements present in the surrounding layer. These results suggest that a reaction between Al, Si, Ti and C took place. Ti 72      Figure 4.19: Line scan data of Ti particle within B319 alloy after Al-5TiCB addition to fluidity mold casting [A]SEM image; [B] Ti line scan; [C] Al line scan; [D] Si line scan A B C D 73  4.3.5 Fluidity Experiments  The fluidity mold developed in this research enabled the comparison of B319 alloy’s fluidity as a function of the type of master alloy additions. The alloy fluidity was related to the mold preheat temperature, the alloy’s pouring temperature and the pressure head driving the flow of the molten B319 aluminum alloy through the mold.  The temperature of the molten B319 aluminum alloy and the mold was measured using K-type thermocouples. The integrated thermocouple at the stopper provided the temperature of the B319 aluminum alloy as the reservoir was being filled until the moment when the plug was pulled and liquid metal entered the mold cavity. On average, the pour time to fill the reservoir was 2-3seconds. During this time period, the thermocouple on the stopper had to increase in temperature from ~400oC to the temperature of the liquid alloy (~705 oC). An example of the thermal data for the liquid B319 aluminum alloy in the reservoir is seen in Figure 4.20.  Figure 4.20:Temperature of liquid metal reservoir in the fluidity mold  When the reservoir was filled and the liquid B319 aluminum alloy had reached the designated metal height, the stopper was pulled out of the reservoir and the liquid B319 aluminum alloy was free to flow into the spiral mold. 300350400450500550600650700750-2 -1 0 1 2 3 4 5Reservior Temperature (C)Time (s)Reservior Temp for Unrefined B319 Fluid Casting First contactAssumed Stopper Removal74  Figure 4.21 shows repeat castings made with unrefined B319 aluminum alloy. The average flow length was 305.5mm, with a standard deviation of 4.5mm.      Figure 4.21: Spiral flow length of unrefined castings [A] Trial#1, [B] Trial#2 A representative thermocouple output of cooling data from three locations along the spiral length (i.e., flow distance in mm) for the unrefined B319 aluminum alloy experiments is seen in Figure 4.22.  Figure 4.22: Cooling curves for unrefined B319 in spiral mold 3003203403603804004204400 5 10 15 20Mold Temperature (C)Time (s)102mm204mm306mmB A Limits of error ±2.2oC 75   This data suggests that B319 aluminum alloy was ~424°C (i.e. partially solid) when it reached a flow distance of 102mm, yet it continued to flow down the spiral to reach a final length of 305.5 mm. As the molten B319 aluminum alloy was fed into the mold, instant solidification occurred at the mold-metal interface. The grain structure of the unrefined B319 sectioned at the 102mm down the spiral’s length is shown in Figure 4.23, and shows a very fine grain structure.    Figure 2.24 shows the SEM microstructure of the unrefined B319 aluminum alloy cross section at the 102mm flow distance.   Figure 4.24: Microstructure (500x) of unrefined B319 sectioned at 102mm from spiral 3mm Figure 4.23: Grain structure of unrefined B319 sectioned at 102mm from spiral Thermocouple Location 76  Figure 4.25: Grain structure of unrefined B319 sectioned at tip of spiral length The microstructure at this flow length revealed fine secondary phase particles evenly distributed within a refined dendritic structure.   Figure 4.25 shows the macrostructure of the unrefined B319 alloy’s tip at the end of the spiral. A grain size gradients (shown with red arrows) from the mold interface toward the center and another from the tip were observed.   Figure 4.26 shows the microstructure of the unrefined B319 aluminum alloy at the frozen metal front.  Figure 4.26: Microstructure of unrefined B319 sectioned tip of spiral length The Al-Mn-Fe intermetallic compounds were seen to coarsen at the frozen metal front. During casting and solidification, the liquid phase in the B319 alloy was mobile and progressed 3mm 77  through solid interdendritic network and likely affected the diffusion field of solutes at the solid-liquid interface [54]. The liquid transported the solutes downstream and increased the concentrations of solute elements at the end (tip) of the spiral, resulting in the coarsening of the secondary phases in this area. Figure 4.27 compares the macrostructure at the spiral tip in fluidity casting of the unrefined B319 alloy and the B319 alloy with master alloy additions.  Unrefined B319   B319 with Al-1TiC Addition    B319 with Al-5TiC Addition   B319 with Al-1TiCB Addition   B319 with Al-5TiCB Addition  Figure 4.27: Macrostructure of B319 alloy at spiral tip. [A] unrefined;  [B] Al-1TiC addition;  [C] Al-5TiC addition; [D] Al-1TiCB addition; [E] Al-5TiCB addition.  3mm A B C D E 78  The addition of the master alloy had no observable changes to the macrostructure of the B319 alloy at the end (tip) of the spiral. The microstructure evaluated with SEM at 2000x magnification in Figure 4.28 below shows no observable structural differences either.   Unrefined B319  B319 with Al-1TiC Addition  B319 with Al-1TiCB Addition  B319 with Al-5TiC Addition  B319 with Al-5TiCB Addition Figure 4.28: Microstructure (2000x) of B319 alloy. [A] unrefined;  [B] Al-1TiC addition;  [C] Al-1TiCB addition; [D] Al-5TiC addition; [E] Al-5TiCB addition. A B C D E 79   Despite the consistent microstructure, Figure 4.29 reveals that the addition of the grain refiner had a statistically significant effect on the B319 alloy’s flow length in the flow spiral.   Figure 4.29: Flow length of all fluidity B319 experiments The results indicate that the addition of master alloys reduced the flow length of the B319 alloy. This result was expected, since the fluidity of aluminum alloys is generally inversely proportional to the alloy’s solidification range [67] (discussed in section 4.3.1). Further, the addition of master alloys had reduced the solidification time of the B319 aluminum alloy. With a constant pressure head maintained throughout all casting experiments, the reduced solidification time would yield a reduced flow length. These results are consistent with the work of Dahle et al. [54] who observed reduced fluidity with Ti addition between 0 - 0.12wt% to an Al-4.5Cu alloy. However, Dahle found that addition of Ti in excess of 0.12wt% increased the fluidity by increasing the fraction of solid at the dendrite coherency point [54].  There were no clearly visible changes in the macro- and microstructure of the B319 alloy in the fluidity mold castings after the addition of the grain refining master alloy. Moreover, there was no correlation between the structure of the spiral tip and the flow length. A possible cause of this insensitivity to grain refinement was the small cross-sectional area of the fluidity spiral, which enabled a high heat extraction from the solidifying alloy. It is well known that the cooling rate during the solidification of aluminum alloys has a great influence on the resulting grain structure. Increased cooling rates allows for reduced undercooling, which  accelerates the rate of nucleation [54]. The B319 alloy refined with master alloy experienced little structural differences 150175200225250275300325Unrefined Al-1TiC Addition Al-5TiC Addition Al-1TiCB Addition Al-5TiCB AdditionFlow Length (mm)80  in the fluidity mold (high cooling rate) compared to the B319 alloy refined with master alloy in the TP-1 mold (low cooling rate). Therefore, in the fluidity spiral castings, the cooling rate had a dominant effect on the grain structure, and superseded that of the inoculating particles. At the mold-metal interface, heat was rapidly extracted resulting in nearly instant solidification. This created a solid “skin”, while a liquid material resided at the core of the casting’s cross section. As the B319 alloy cooled, the forming dendrites interacted with neighboring dendrites. Eventually the dendrites interlocked and began to resist further flow. However, incoming heat may have partially re-melted and/or fractured weaker dendrites that were forming on the skin and carried the dendrite fragments downstream, thereby promoting heterogenous nucleation. This was observed in the work of Dahle et al. [54] who tested the effect of grain refinement on the fluidity of foundry alloys using a sand spiral mold. The nature of the spiral fluidity experiment inherently resulted in the transport of solid dendrites downstream. The combination of free crystals (α-Al solids and solute particles) and high cooling rates jointly contributed to a fine equiaxed structure in the unrefined B319 alloy.  The flow length of the B319 alloy refined with the Al-1TiCB and Al-5TiCB master alloys was comparable and within the repeatability measurement error. No observable differences in macro and microstructure was recorded between the Al-1TiCB and Al-5TiCB experiments; however, (as mentioned in section 4.3.4) Ti particles were found within the microstructure of the cross section of B19 with Al-5TiCB (at 102mm along the flow length). Figure 4.30 shows the SEM/EDS image with a group of Ti particles.    81    The addition of Ti to pure Al has been reported to increase the viscosity of the B319 alloy [50], [68].  Yu et al. and Banerji et al. [50], [69] reported that the increased viscosity could be a result of two effects. The first being the adding solid particles to a melt; and the second being a change of the melt structure. The Ti particles shown in Figure 4.30 (and previously discussed in section 4.3.4) have likely reacted with the surrounding Al, and therefore not only changed the physical properties of the interface between the Al and the coated Ti particle, but also likely changed the chemistry of the intergranular liquid.    Figure 4.30: SEM-EDS image of Ti cluster in B319 casting with Al-5TiCB addition.[A]SEM/EDS image; [B] Al; [C] Si; [D]; Ti; [E] Fe  A B C D E 82  4.3.6 Mechanical Properties  The mechanical properties of the as-cast materials were evaluated using tensile and hardness testing. The tensile samples and procedure were in accordance with the ASTM B557M standard. The hardness testing was done in accordance with the ASTM E18 standard. For each tensile experiment, two samples were extracted from the TP1 casting. All hardness reading were taken from the samples used to evaluate the microstructure of TP1 castings.  4.3.6.1 Tensile Properties Figure 4.31 shows the results of ultimate tensile strength (UTS) measurement for all the as-cast alloys.  The B319 alloy with Al-5TiC master alloy and B319 alloy with Al-5TiCB master alloy had a higher UTS than the unrefined B319 alloy by 6% and 8%, respectively. This was expected due to the reduced grain size and a homogeneous distribution of fine particles in the microstructure.   Figure 4.31: UTS and elongation of refined and unrefined B319 In contrast, the B319 alloy with Al-1TiC and Al-1TiCB master alloy additions had a lower UTS than the unrefined B319 alloy by 10% and 14%, respectively. This result was not consistent with the reduced grain size discussed in section 4.3.2. The exact mechanism for these results remains unclear. 0123456200220240260280300320340Unrefined Al-5TiC Addition Al-5TiCB Addition Al-1TiC Addition Al-1TiCB AdditionElonggation (%)Ultimate Tensile Strength (MPa)UTS Elongation83  In general, the grain size and elongation were expected to be inversely proportional. However, the elongation remained constant after at the addition Al-5TiC and Al-5TiCB to the B319 alloy while the addition of Al-1TiC and Al-1TiCB both reduced the elongation of the as-cast alloy. Kumar and Bichler have reported that unrefined B319 alloys exhibit brittle fracture with samples experiencing an elongation of 1-2% [34]. In their study, brittle secondary phases (CuAl2 and Fe intermetallics) were responsible for acting as stress concentrations that promoted brittle fracture under axial loading. Because the B319 alloy is naturally brittle, the effect of the grain refiners on the alloy’s elongation is difficult to investigate in as-cast condition.  The yield  strength and Young’s modulus from the grain refining experiments are shown in Figure 4.32.  Figure 4.32: Yield strength and Young’s modulus of refined and unrefined B319 All grain refining experiments did not have statistically significant impact on the yield strength; however, the Al-1TiC and Al-1TiCB refined alloys exhibited a minor decrease in Young’s modulus. It is well known that fine intermetallics that are evenly dispersed will increase the likelihood that they interfere with the slip process. Therefore, the reduction in young’s modulous is the result of coarsening of intermetallics (discussed in section 4.3.3).  0246810121416100110120130140150160170Unrefined Al-5TiC Addition Al-5TiCB Addition Al-1TiC Addition Al-1TiCB AdditionYoungs Modulous  (GPa)Yeild Strength (MPa)Yeild Strength Young's Modulus84  4.3.6.2 Hardness Figure 4.33 shows the effect of the master alloys on the hardness of the B319 aluminum alloy. The addition of the grain refiners to the B319 aluminum alloy was seen to refine grains and result in smaller secondary phases. Hard intermetallic phases, such as the Cu and Fe-rich intermetallics, are generally known to influence the hardness of the as-cast alloy [27], [36]. The addition of all master alloys increased the hardness of the B319 alloy. The addition of pure Al had no significant effect on the hardness of the B319, therefore varying the Al concentration by 3wt% did not result in any appreciable hardness variation. The Powder TiC addition resulted in a small increase in the hardness; however, the change was not observed to be as large as reported in the work of Kumar [34]. Reasons for this discrepancy may be a result of different mixing and pouring techniques.    Figure 4.33: Hardness of refined and unrefined B319   6062646668707274767880Unrefined Pure AlAdditionTiC PowderAdditionAl-1TiCAdditionAl-5TiCAdditionAl-1TiCBAdditionAl-5TiCBAdditionHR15T85  Chapter 5: Conclusions This chapter reviews the highlights of this thesis work. The conclusions are categorized into three categories: 1) SPS Master Alloy Characterization; 2) B319 Aluminum Alloy Characterization; 3) Master Alloy Experiments 5.1 SPS Master Alloy Characterization The SPS Master Alloys were characterized for their particle distribution, structural features, and densification. The following summarizes the findings of this thesis: • The SPS process successfully produced a TiC and TiCB Master Alloy with ~2-4% porosity using Al as the base material. • The physical segregation of the fine TiC on the larger Al particles was the result of the significant particle size difference. 5.2 B319 Aluminum Alloy Characterization The B319 alloy was characterized for microstructure features, grain size, and thermal profile. The following lists the findings of the characterization of the unrefined B319 alloy studied in this thesis: • The B319 alloy consists of an aluminum matrix solid solution with secondary eutectic phases and intermetallic particles forming between the primary Al grains. • The Si-rich phases precipitating in the B319 alloy were nearly pure silicon and took on an acicular morphology • The dominant secondary phase in the B319 aluminum alloy was the Al-Mn-Fe intermetallic phase, with a large script-like morphology. • B319 alloy contained a copper-rich intermetallic, with a stoichiometry of CuAl2 and block-like morphology. • The grain structure of the unrefined as-cast B319 aluminum alloy had an equiaxed structure due to solute elements providing constitutional undercooling.  86  5.3 Master Alloy Experiments • All experiments with SPS master alloy addition eliminated the undercooling that was observed in the unrefined B319 aluminum alloy. • When the Al-5TiC and Al-5TiCB was added, the α-Al nucleation temperature (TN,α) increased from 654.4°C  to 659.3°C  and 659.9°C , respectively.  • The addition of Al-1TiC and Al-1TiCB when the α-Al nucleation temperature increased from 654.4 °C to 660.5°C  and 659.1°C , respectively. • Al-5TiC and Al-5TiCB increased the total cooling rate of the B319 alloy by increasing the solidification range and reducing the solidification time due to increased numbers of growing solids during solidification. • The silicon eutectic nucleation temperature (TN,Si) on average increased 5 °C after the addition of all master alloys. The copper eutectic nucleation temperature (TN,Cu) on average increased 7 °C after the addition of all master alloys. • Addition of the Al-1TiC and Al-5TiC master alloys yielded a homogeneous distribution of the intermetallic and eutectic particles throughout the as-cast material.  • All SPS master alloys (Al-1TiC, Al5TiC, Al-1TiCB, Al5TiCB) reduced the grain size to approximately 40% of the size of the unrefined B319 aluminum alloy.  • The flow length of the B319 was reduced by 15%, 31%, 29%, and 32% after addition of Al-1TiC, Al-5TiC, Al-1TiCB  and Al-5TiCB, respectively. • The B319 alloy with Al-5TiC master alloy and B319 alloy with Al-5TiCB master alloy had a higher UTS than the unrefined B319 alloy by 6% and 8%, respectively. • The B319 alloy with Al-1TiC master alloy and B319 alloy with Al-1TiCB master alloy had a lower UTS than the unrefined B319 alloy by 10% and 14%, respectively • The hardness of the B319 was increased by 10%, 11%, 11%, and 13% after addition of Al-1TiC, Al-5TiC, Al-1TiCB  and Al-5TiCB, respectively.   87  Chapter 6: Future Work This chapter provides suggestions for future work to further understand the solidification characteristics and behavior of the B319 aluminum alloy 6.1 SPS Master Alloy Some recommendations for further investigation and testing of the TiC-Based master alloys via SPS: • Test the effect of TiC particle size on the solidification behavior and resulting microstructure of B319. • Develop varying high concentration SPS master alloy pellets to determine if TiC particle agglomeration occurs more frequently • Investigate the effect of a CB-based master alloy to the B319 alloy 6.2 Fluidity • Modify spiral fluidity mold used in this study to see if results are repeatable. Design a new style of testing apparatus to better control the melt temperature by inserting a quartz tubing into the crucible and pulling the molten material with attached pump.    88  Bibliography  [1] “Environment and Climate Change Canada - Climate Change - Climate Change Science and Research,” 2015. [Online]. Available: http://www.ec.gc.ca/sc-cs/. [Accessed: 06-Sep-2016]. [2] O. US EPA, “Sources of Greenhouse Gas Emissions,” 2016. [Online]. Available: https://www.epa.gov/ghgemissions/sources-greenhouse-gas-emissions. [Accessed: 06-Sep-2016]. [3] “Top 10 Factors Contributing To Fuel Economy - Car Maintenance and Car Repairs - DriverSide,” 2015. [Online]. Available: http://www.driverside.com/auto-library/top_10_factors_contributing_to_fuel_economy-317. [Accessed: 06-Oct-2016]. [4] “Vehicle Technologies Office: Lightweight Materials for Cars and Trucks | Department of Energy,” 2014. [Online]. Available: http://energy.gov/eere/vehicles/vehicle-technologies-office-lightweight-materials-cars-and-trucks. [Accessed: 06-Oct-2016]. [5] A. D. Brooker, L. Wang, J. Ward, and D. of Energy, “Lightweighting Impacts on Fuel Economy, Cost, and Component Losses,” SAE Int., pp. 1–10, 2013. [6] Z. Li, A. . Samuel, F. . Samuel, C. Ravindran, S. Valtierra, and H. . Doty, “Parameters controlling the performance of AA319-type alloys: Part I. Tensile properties,” Mater. Sci. Eng. A, vol. 367, no. 1, pp. 96–110, 2004. [7] J. R. Davis, “Aluminum and Aluminum Alloys,” in Alloying: Understanding the Basics, 2001: ASM International, 2001, pp. 351–416. [8] J. G. Kaufman and E. L. Rooy, Aluminum Alloy Castings: Properties , Processes , and Applications. 2004. [9] H. Nguyen, “Manufacturing Processes and Engineering Materials Used in Automotive Engine Blocks,” Grand Valley University, 2005. [10] A. M. Samuel and F. H. Samuel, “Effect of melt treatment, solidification conditions and porosity level on the tensile properties of 319.2 endchill aluminium castings,” J. Mater. Sci., vol. 30, no. 19, pp. 4823–4833, 1995.   89   [11] F. Farhang Mehr, C. Reilly, S. Cockcroft, D. Maijer, and R. MacKay, “Effect of chill cooling conditions on cooling rate, microstructure and casting/chill interfacial heat transfer coefficient for sand cast A319 alloy,” Int. J. Cast Met. Res., vol. 27, no. 5, pp. 288–300, Oct. 2014. [12] J. E. Gruzleski, Microstructure Development during Metal Casting. Montreal, Canada: American Foundrymen’s Society, Inc., 2000. [13] G. K. Sigworth and T. A. Kuhn, “Grain refinement of aluminum casting alloys,” Int. J. Met., vol. 1, no. 1, pp. 31–40, 2007. [14] T. E. Quested, “Understanding mechanisms of grain refinement of aluminium alloys by inoculation Understanding mechanisms of grain refinement of aluminium alloys by inoculation Grain refining practice Introduction to grain refinement,” Mater. Sci. Technol., vol. 20, 2004. [15] Z. Li, A. . Samuel, F. . Samuel, C. Ravindran, H. . Doty, and S. Valtierra, “Parameters controlling the performance of AA319-type alloys: Part II. Impact properties and fractography,” Mater. Sci. Eng. A, vol. 367, no. 1, pp. 111–122, 2004. [16] A. K. P. Rao, B. S. Murty, M. Chakraborty, and M. Chakraborty, “Improvement in tensile strength and load bearing capacity during dry wear of Al–7Si alloy by combined grain refinement and modification,” Mater. Sci. Eng. A, vol. 395, no. 1–2, pp. 323–326, Mar. 2005. [17] N. Hansen, “Hall–Petch relation and boundary strengthening,” Scr. Mater., pp. 801–806, 2004. [18] T. E. Quested and A. L. Greer, “Grain refinement of Al alloys: Mechanisms determining as-cast grain size in directional solidification,” Acta Mater., vol. 53, no. 17, pp. 4643–4653, 2005. [19] P. S. Mohanty and J. E. Gruzleski, “Grain refinement mechanisms of hypoeutectic Al-Si alloys,” Acta Mater., vol. 44, no. 9, pp. 3749–3760, 1996. [20] Z. Fan et al., “Grain refining mechanism in the Al/Al–Ti–B system,” Acta Mater., vol. 84, 90  pp. 292–304, 2015. [21] “A Guide to Aluminum Casting Alloys,” 2005. [Online]. Available: http://www.mid-atlanticcasting.com/. [Accessed: 18-Oct-2016]. [22] S. K. Chaudhury, D. Apelian, P. Meyer, D. Massinon, and J. Morichon, “Microstructure and Mechanical Properties of Heat-Treated B319 Alloy Diesel Cylinder Heads,” Metall. Mater. Trans. A, vol. 46, no. 7, pp. 3276–3286, Jul. 2015. [23] Q. L. Bai, Y. Li, H. X. Li, Q. Du, J. S. Zhang, and L. Z. Zhuang, “Roles of Alloy Composition and Grain Refinement on Hot Tearing Susceptibility of 7××× Aluminum Alloys,” Metall. Mater. Trans. A, vol. 47, no. 8, pp. 4080–4091, 2016. [24] M. G. Kalhapure and P. M. Dighe, “Impact of Silicon Content on Mechanical Properties of Aluminum Alloys,” Int. J. Sci. Res. ISSN (Online Index Copernicus Value Impact Factor, vol. 14, no. 6, pp. 2319–7064, 2013. [25] A. M. Samuel, H. W. Doty, S. Valtierra, and F. H. Samuel, “Defects related to incipient melting in Al–Si–Cu–Mg alloys,” Mater. Des., vol. 52, pp. 947–956, 2013. [26] P. Marchwica, J. H. Sokolowski, and W. T. Kierkus, “Fraction solid evolution characteristics of AlSiCu alloys -dynamic baseline approach,” AMMA, vol. 47, no. 2, pp. 115–136, 2001. [27] M. Tash, F. H. Samuel, F. Mucciardi, and H. W. Doty, “Effect of metallurgical parameters on the hardness and microstructural characterization of as-cast and heat-treated 356 and 319 aluminum alloys,” Mater. Sci. Eng. A, vol. 443, no. 1–2, pp. 185–201, 2007. [28] S. Ji, D. Watson, Y. Wang, M. White, and Z. Fan, “Effect of Ti Addition on Mechanical Properties of High Pressure Die Cast Al-Mg-Si Alloys,” Trans Tech Publ., pp. 23–27, 2013. [29] M. Nowak and H. B. Nadendla, “Grain Refiner for Aluminium-Silicon Sand Casting Aloys,” in Light Metals 2012, Hoboken, NJ, USA: John Wiley & Sons, Inc., 2012, pp. 349–353. [30] J. A. Taylor, G. B. Schaffer, and D. H. StJohn, “The role of iron in the formation of porosity in Al-Si-Cu-based casting alloys: Part III. A microstructural model,” Metall. 91  Mater. Trans. A, vol. 30, no. 6, pp. 1657–1662, Jun. 1999. [31] Y. Jinku, Q. Qi, L. Lian, J. Qihua, and N. Dongying, “Investigation on the undercooling and crystallization of pure aluminum melt by DSC,” Phase Transitions, vol. 83, no. 7, pp. 543–549, 2010. [32] R. Barua, “Kinetics of Phase Transformations: Nucleation and Growth,” Northeastern Univeristy. [Online]. Available: https://pdfs.semanticscholar.org/presentation/3a1c/c61fc854c76c41cb7c9672740f88d04f0e32.pdf. [Accessed: 07-Oct-2016]. [33] M. Duffy, “Nucleation - Soft-Matter,” 2011. [Online]. Available: http://soft-matter.seas.harvard.edu/index.php/Nucleation. [Accessed: 15-Dec-2016]. [34] V. Kumar and L. Bichler, “Effect of TiC Addition on the Microstructure and Mechanical Properties of B319 Alloy,” Trans. Indian Inst. Met., vol. 68, no. 6, pp. 1173–1180, Dec. 2015. [35] A. K. P. Rao, K. Das, B. S. Murty, and M. Chakraborty, “Al–Ti–C–Sr master alloy—A melt inoculant for simultaneous grain refinement and modification of hypoeutectic Al–Si alloys,” J. Alloys Compd., vol. 480, no. 2, pp. L49–L51, Jul. 2009. [36] T. Wang, T. Gao, P. Zhang, J. Nie, and X. Liu, “Influence of a new kind of Al–Ti–C master alloy on the microstructure and mechanical properties of Al–5Cu alloy,” J. Alloys Compd., vol. 589, pp. 19–24, 2014. [37] S. Nakane, Y. Takano, M. Yoshinaka, and K. Hirota, “Fabrication and Mechanical Properties of Titanium Boride Ceramics,” vol. 28, pp. 1627–1628, 1999. [38] X. Wang, Z. Liu, W. Dai, and Q. Han, “On the Understanding of Aluminum Grain Refinement by Al-Ti-B Type Master Alloys,” Metall. Mater. Trans. B, vol. 46, no. 4, pp. 1620–1625, Aug. 2015. [39] G. E. Dieter and D. Bacon, METALLURGY SI Metric Edition Adapted by. . [40] S. G. Shabestari and M. Malekan, “Assessment of the effect of grain refinement on the solidification characteristics of 319 aluminum alloy using thermal analysis,” J. Alloys Compd., vol. 492, no. 1–2, pp. 134–142, 2010. 92  [41] T. Wang, T. Gao, J. Nie, P. Li, and X. Liu, “Influence of carbon source on the microstructure of Al–Ti–C master alloy and its grain refining efficiency,” Mater. Charact., vol. 83, pp. 13–20, 2013. [42] B. Q. Zhang, H. S. Fang, L. Lu, M. O. Lai, H. T. Ma, and J. G. Li, “Synthesis Mechanism of an Al-Ti-C Grain Refiner Master Alloy Prepared by a New Method.” [43] “Manufacturing Process of Carbon Black.” [Online]. Available: http://www.carbonblack.jp/en/cb/seizou.html. [Accessed: 07-Oct-2017]. [44] J.-B. Donnet, R. C. Bansal, and M.-J. Wang, Carbon black: Science and Technology. Marcel Dekker, 1993. [45] I.-S. Ahn, S.-Y. Bae, H.-J. Cho, C.-J. Kim, and D.-K. Park, “Synthesis of Titanium Carbide by Heat Treatment of TiH 2 and Carbon Black Powders,” Trans Tech Publ., vol. 534–536, pp. 255–228, 2007. [46] M. Razavi, M. R. Rahimipour, and R. Kaboli, “Synthesis of TiC nanocomposite powder from impure TiO2 and carbon black by mechanically activated sintering,” J. Alloys Compd., vol. 460, no. 1–2, pp. 694–698, 2008. [47] J. B. Holt and Z. A. Munir, “Combustion synthesis of titanium carbide: Theory and experiment,” J. Mater. Sci., vol. 21, no. 1, pp. 251–259, Jan. 1986. [48] A. Banerji and W. Reif, “Grain refinement of aluminum by TiC,” Metall. Trans. A, vol. 16, no. 11, pp. 2065–2068, 1985. [49] P. S. Mohanty and J. E. Gruzleski, “GRAIN REFINEMENT OF ALUMINIUM BY TiC,” vol. 31, no. 2, pp. 179–184, 1994. [50] A. Banerji and W. Reif, “Development of Al-Ti-C grain refiners containing TiC,” Metall. Mater. Trans. A, vol. 17, no. 12, pp. 2127–2137, 1986. [51] A. Banerji, W. Reif, and Q. Feng, “Metallographic investigation of TiC nucleants in the newly developed AI-Ti-C grain refiner,” J. Mater. Sci., vol. 29, pp. 1958–1965, 1994. [52] C. A. Leon, V. H. Lopez, E. Bedolla, and R. A. L. Drew, “Wettability of TiC by commercial aluminum alloys,” J. Mater. Sci., vol. 37, no. 16, pp. 3509–3514, 2002. [53] N. Froumin, N. Frage, M. Polak, and M. P. Dariel, “WETTING PHENOMENA IN THE 93  TiC/(Cu±Al) SYSTEM,” Acta Mater., vol. 48, pp. 1435–1441, 2000. [54] A. K. Dahle, P. A. Tøndel, C. J. Paradies, and L. Arnberg, “Effect of grain refinement on the fluidity of two commercial Al-Si foundry alloys,” Metall. Mater. Trans. A, vol. 27, no. 8, pp. 2305–2313, Aug. 1996. [55] N. L. M. Veldman, A. K. Dahle, D. H. StJohn, and L. Arnberg, “Dendrite coherency of Al-Si-Cu alloys,” Metall. Mater. Trans. A, vol. 32, no. 1, pp. 147–155, Jan. 2001. [56] L. Bäckerud, G. Chai, and J. Tamminen, Solidification characteristics of aluminium alloys - Foundry Alloys, vol. 2. 1990. [57] A. K. Prasada Rao, K. Das, B. S. Murty, and M. Chakraborty, “Microstructural and wear behavior of hypoeutectic Al–Si alloy (LM25) grain refined and modified with Al–Ti–C–Sr master alloy,” Wear, vol. 261, no. 2, pp. 133–139, Jul. 2006. [58] G. Xie, “Powder Metallurgy &amp; Mining Spark Plasma Sintering: A Useful Technique to Develop Large-Sized Bulk Metallic Glasses,” Xie. J Powder Met. Min, vol. 2, no. 2, 2013. [59] G. A. Sweet, M. Brochu, R. L. Hexemer, I. W. Donaldson, and D. P. Bishop, “Effect of Transition Metal Additions on the Spark Plasma Sintering Response of Aluminum Powders,” Powder Metall. World Congr., no. August, 2014. [60] Z. Trzaska, A. Couret, and J.-P. Monchoux, “Spark plasma sintering mechanisms at the necks between TiAl powder particles,” 2016. [61] D. V Lazurenko, V. I. Mali, N. S. Belousova, and A. Theommes, “Formation of Intermetallic Structures by Spark Plasma Sintering of Titanium and Aluminum Powders,” Trans Tech Publ., pp. 179–181, 2015. [62] ASTM B557-M, “Standard Test Methods for Tension Testing Wrought and Cast Aluminum- and Magnesium-Alloy Products (Metric)1,” ASTM, vol. 8, no. 4, 2015. [63] D. Culliton, T. Betts, D. Kennedy, D. Culliton, A. J. Betts, and D. Kennedy, “Impact of Intermetallic Precipitates on the Tribological and/or Corrosion Performance of Cast Aluminium Alloys: a Short Review,” Int. J. Cast Met. Res., vol. 26, no. 2, 2013.  94  [64] L. Svendsen and A. Jarfors, “Al–Ti–C phase diagram,” Mater. Sci. Technol., vol. 9, no. 11, pp. 948–957, 1993. [65] S. G. Shabestari and M. Malekan, “Thermal Analysis Study of the Effect of the Cooling Rate on the Microstructure and Solidification Parameters of 319 Aluminum Alloy,” vol. 44, no. 3, pp. 305–312, 2005. [66] A. Flores-Valdes, M. . Rivas-Aguilar, and J. . Escobedo-Bocardo, “Application of a modified temperature difference thermal analysis technique to study the solidification of an Al Si 319 alloy,” Mater. Des., vol. 14, no. 4, pp. 223–229, 1993. [67] K. R. Ravi, R. M. Pillai, K. R. Amaranathan, B. C. Pai, and M. Chakraborty, “Fluidity of aluminum alloys and composites: A review,” J. Alloys Compd., vol. 456, no. 1–2, pp. 201–210, 2008. [68] L. F. (Lucio F. . Mondolfo, Aluminum alloys : structure and properties. Butterworths, 1979. [69] L. Yu and X. Liu, “Ti transition zone on the interface between TiC and aluminum melt and its influence on melt viscosity,” J. Mater. Process. Technol., vol. 182, no. 1–3, pp. 519–524, 2007. [70] W. D. Callister and D. G. Rethwisch, Materials Science and Engineering An Introduction. John Wiley & Sons, Inc., 2007.    95  Appendix  𝑇ℎ𝑒𝑜𝑟𝑒𝑡𝑖𝑐𝑎𝑙 𝐷𝑒𝑛𝑠𝑖𝑡𝑦 𝑜𝑓 𝑆𝑃𝑆 𝐵𝑙𝑒𝑛𝑑𝑠= [(𝑤𝑒𝑖𝑔ℎ𝑡 % 𝑜𝑓 𝑚𝑎𝑡𝑒𝑟𝑖𝑎𝑙 𝐴) × (𝑑𝑒𝑛𝑠𝑖𝑡𝑦 𝑜𝑓 𝑚𝑎𝑡𝑒𝑟𝑖𝑎𝑙 𝐴)]+ [(𝑤𝑒𝑖𝑔ℎ𝑡 % 𝑜𝑓 𝑚𝑎𝑡𝑒𝑟𝑖𝑎𝑙 𝐵) × (𝑑𝑒𝑛𝑠𝑖𝑡𝑦 𝑜𝑓 𝑚𝑎𝑡𝑒𝑟𝑖𝑎𝑙 𝐵)] Sample Calculation: Al-5TiC: • 95wt% Aluminum at 2.70g/cm3 • 5wt% Titanium Carbide at 4.93g/cm3 𝐴𝑙 − 5𝑇𝑖𝐶 𝑇ℎ𝑒𝑜𝑟𝑒𝑡𝑖𝑐𝑎𝑙 𝐷𝑒𝑛𝑠𝑖𝑡𝑦 = [(0.95) × (2.7)] + [(0.05) × (4.93)] =  𝟐. 𝟖𝟏𝟐𝒈/𝒄𝒎𝟑 Al-5TiCB • 95wt% Aluminum at 2.70g/cm3 • 4.6wt% Titanium at 4.506g/cm3 • 0.4wt% Carbon black at 2.0g/cm3 NOTE: Carbon black has amorphous and has a density range of 1.8-2.1g/cm3 𝐴𝑙 − 5𝑇𝑖𝐶𝐵 𝑇ℎ𝑒𝑜𝑟𝑒𝑡𝑖𝑐𝑎𝑙 𝐷𝑒𝑛𝑠𝑖𝑡𝑦 = [(0.95) × (2.7)] + [(0.046) × (4.506)] + [(0.004) × (2.0)]=  𝟐. 𝟕𝟕𝟐𝒈/𝒄𝒎𝟑         

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