A Study of Fluidity, Flow Stoppage and Silicon Modification in Lost Foam Casting of Hypereutectic Aluminum-Silicon Alloys

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A STUDY OF FLUIDITY, FLOW STOPPAGE AND SILICON MODIFICATION IN LOST FOAM CASTING OF HYPEREUTECTIC ALUMINUM-SILICON ALLOYS

by Achinta Haldar B.Sc. (Engg.), Metallurgical Engineering (1990) National Institute of Technology, Rourkela, India

A thesis presented to Ryerson University

in partial fulfillment of the requirements for the degree of Master of Applied Science in the Program of Mechanical Engineering

Toronto, Ontario, Canada, 2007 © Achinta Haldar 2007

ISBN: 978-0-494-41374-6

AUTHOR’S DECLARATION

I hereby declare that I am the sole author of this thesis or dissertation. I authorize Ryerson University to lend this thesis or dissertation to other institutions or individuals for the purpose of scholarly research.

Achinta Haldar

I further authorize Ryerson University to reproduce this thesis or dissertation by photocopying or by other means, in total or in part, at the request of other institutions or individuals for the purpose of scholarly research.

Achinta Haldar

ii BORROWER’S PAGE

Name Address Date (last, first) (day/month/year)

iii ABSTRACT

A STUDY OF FLUIDITY, FLOW STOPPAGE AND SILICON MODIFICATION IN LOST FOAM CASTING OF HYPEREUTECTIC ALUMINUM-SILICON ALLOYS

Hypereutectic Al-Si alloys were cast by the lost foam casting (LFC) process to study the effect of silicon level and superheat on the alloy fluidity and silicon morphology. Three alloys with 14% Si, 18% Si and 22%Si were cast at three superheats 60ºC, 115ºC and 170ºC. Expendable poly styrene (EPS) foam spirals, 12.50 mm wide and with two different thicknesses, 6.25 mm and 12.50 mm were used. It was observed that metal fluidity increased with increases in silicon level and superheat. The investigation revealed that the interlocked primary silicon crystals (PSC) agglomerated behind the flow tip, which contained the eutectic silicon. The distance between the flow tip and the PSC agglomerated zone increased with an increase in silicon level. Significant PSC floatation was observed for Al-22%Si alloy. The size of the silicon depleted zone increased with an increase in superheat. However, this trend was reduced by the addition of phosphorus, which modified the shape and reduced the size of PSC. Increasing the ceramic master alloy (CMA) content in the alloy resulted in a progressive modification of the eutectic silicon from needle-like to lamellar and finally to fibrous shape. Further, the addition of CMA induced precipitation of Mg rich phases. There was an increase in the surface hardness of the alloys due to addition of CMA as well as phosphorus.

iv ACKNOWLEDGEMENTS

First of all I wish to express my gratitude to Professor C. Ravindran for giving me the unique opportunity to complete this thesis. His constant advice, support and encouragement made this dream come true. I remain highly indebted to this great man. I take this opportunity to thank A. Machin (Technical Officer) for his external help, advice and support during the entire program. His ideas were simply great. I also appreciate J. Amankrah (Technical Officer) and Q. Li (Technical Officer) for their support during the experiments and analysis. I appreciate L. Bichler for his help during preparation of the thesis. S. Jagoo, K. Lee, A. Elsayed and F. D’Alia are also acknowledged for their help during the experimental work. I am very grateful to Dr. J. Sokolowski and his research team at the University of Windsor for allowing me to use their facility and equipment to derive very important data. I am very thankful to D. Nolan of Foseco Morval, Ontario, for his advice and providing the required materials for experiments. I also thank L. Pike of Gamma Foundries, Ontario for his help in doing the chemical analysis of my castings. I heartily acknowledge AUTO21 network of excellence for funding and supporting the entire project. I also thank Dr. T. Troczynski of University of British Columbia for supplying the ceramic master alloy and Linamar Corporation for supplying the A319 ingots. Special thanks to my colleague, N. Dahata for his contributions during building of the rotary argon degasser at the lab. His continuous help and support during the entire period of study is also appreciated. I also thank my friend S. Gupta for his moral support and assistance during compilation of my thesis draft. My wife Mahuya had been my constant support and inspiration during all the ups and downs during the period. I remain indebted to her understanding and help especially during the latter period. I am lucky to have a little son Mayukh who too had been very supportive. My love to my little angels Amogha and Adrika. I also pay my respects to my parents for their constant selfless prayers for my well- being and fulfillment of my dreams. What I am today is because for them. Finally, I thank GOD for everything that I have achieved till now and pray for what I still desire to achieve.

v TABLE OF CONTENTS

TITLE PAGE i AUTHOR’S DECLARATION ii BORROWER'S PAGE iii ABSTRACT iv ACKNOWLEDGEMENTS v TABLE OF CONTENTS vi LIST OF TABLES x LIST OF FIGURES xii NOMENCLATURE xvi

Chapter 1. INTRODUCTION 1 Research objective 2

Chapter 2. LITERATURE REVIEW 3 2.1. Aluminum alloys 3 2.1.1. Aluminum Silicon phase diagram 4 2.1.2. Aluminum Silicon casting alloys 4 2.2. Lost Foam Casting (LFC) process 5 2.3. Fluidity Study 6 2.3.1. Flow length 6 2.3.2. Metal velocity 8 2.4. Mold feeding and solidification 8 2.4.1. Feeding mechanism 8 2.4.2. Castability and solidification characteristics of Al-Si alloys 10 2.5. Flow stoppage mechanism in metals 11 2.5.1. Flow stoppage mechanism for pure metals 12 2.5.2. Flow stoppage mechanism for hypoeutectic alloys 12 2.5.3. Flow stoppage mechanism for hypereutectic Al-Si alloys 13

vi 2.6. Thermal analysis 15 2.6.1. Cooling curve analysis 15 2.6.2. Solidification time and range 16 2.7. Precipitation reactions during solidification of Al-Si alloys 17 2.7.1. Hypoeutectic A319 alloy 17 2.7.2. Hypereutectic A390 alloy 18 2.8. Silicon phases 19 2.8.1. Eutectic and PSC morphology 19 2.8.2. Floatation of PSC 20 2.9. Silicon modification 20 2.9.1. Eutectic silicon modification mechanism 22 2.9.2. Primary silicon modification mechanism 22 2.9.3. Effects on mechanical properties 22 2.9.4. Types of silicon modifiers 23 2.9.4.1. Ceramic Master Alloy (CMA) 23 2.9.4.2. Phosphorus 24 2.10. Scanning electron microscopy (SEM) and 26 X-Ray energy dispersive spectroscopy (XEDS)

Chapter 3. EXPERIMENTAL PROCEDURE 27 3.1. Research plan 27 3.2. Materials 29 3.2.1. Al-Si alloys 29 3.2.2. Silicon modifiers 30 3.2.3. EPS foam boards 31 3.2.4. Coating 32 3.2.5. Mold media 33 3.3. Pattern making 33 3.3.1. Coating and drying 36

vii 3.4. Experiment setup 37 3.4.1. Preparation and position of thermocouples 37 3.4.2. Pattern embedding and sand compaction 39 3.4.3. Connection to data acquisition 40 3.5. Melting and alloy preparation 40 3.5.1. Degassing and measurement of hydrogen content 42 3.5.2. Addition of modifiers 44 3.6. Pouring procedure 44 3.6. Analysis of castings 45 3.7.1. Measurement of fluidity 45 3.7.2. Thermal data analysis 45 3.7.3. UMSA 46 3.7.4. Sampling for analysis 48 3.7.5. Hardness testing 48 3.7.6. Sample preparation for Image analysis 48 3.7.7. Measurement of PSC size 50 3.7.8. SEM analysis 50

Chapter 4. RESULTS AND DISCUSSION 51 4.1. Flow characteristics of unmodified Al-Si alloys 51 4.1.1. Flow length 51 4.1.2. Metal velocity 53 4.1.3. Study of flow tip and gate area 54 4.1.3.1. Flow tips 56 4.1.3.2. Gate area 58 4.2. Silicon morphology in unmodified Al-Si alloys 59 4.2.1. Eutectic silicon 59 4.2.2. Primary silicon 60 4.3. The PSC floatation phenomenon 63 4.3.1. PSC distribution across the cross-section height 64 4.3.2. Size and distribution of PSC 65

viii 4.4. Effect of PSC on flow stoppage 66 4.5. Effect of CMA on fluidity and alloy microstructure 68 4.5.1. Flow length 68 4.5.2. Eutectic silicon morphology 69 4.5.3. PSC morphology 71 4.5.4. PSC floatation 71 4.6. Effect of phosphorus on fluidity and alloy microstructure 72 4.6.1. Flow length 72 4.6.2. Eutectic silicon morphology 73 4.6.3. PSC morphology 73 4.6.4. PSC floatation 75 4.7. Surface hardness studies 76 4.7.1. Micro hardness tests 76 4.7.2. Surface hardness of the unmodified alloys 77 4.7.3. Effects of modifiers on surface hardness 78 4.7.3.1. Effects of CMA 78 4.7.3.2. Effects of phosphorus 79 4.8. Thermal analysis 82 4.9. Precipitated phases and intermetallic compounds 84 4.9.1. X-Ray mapping and XEDS 85 4.9.2. Identified Phases 86 4.10. Combined effect of phosphorus and CMA 89

Chapter 5. CONCLUSIONS AND SUGGESTIONS FOR FURTHER STUDY 90 5.1. Effect of silicon and superheat on alloy fluidity 90 5.2. Flow stoppage mechanism 91 5.3. Effect of modifiers on eutectic silicon and PSC morphology 91

REFERENCES 93

APPENDIX 101

ix LIST OF TABLES

Description Page# Table 2.1. Stages of solidification and their effects 11 Table 2.2. Composition of typical A390 alloy 11 Table 2.3. Summary of reaction during solidification of A319 alloy 18 Table 2.4. Summary of reaction during solidification of A390 alloy 19 Table 2.5 PSC size in Al-22%Si alloy treated with 0.055%P 25 Table 3.1 Composition of A319 alloy used for the experiment 29 Table 3.2. Composition of the Al-50%Si hardener alloy 30 Table 3.3. Ceramic master alloy (CMA) composition 30 Table 3.4. Ceramic master alloy (CMA) general information 31 Table 3.5. Fos Flo-7, copper-phosphorus brazing alloy composition 31 Table 3.6. Specifications of the EPS foam 32 Table 3.7. Specifications of the coating 32 Table 3.8. Dimensions of the spirals 33 Table 3.9.(a). Position of thermocouples on 12.50 mm thick spiral 39 Table 3.9.(b). Position of thermocouples on 6.25 mm thick spiral 39 Table 3.10. Sequence of activities and temperature schedules 41 Table 3.11. Pouring temperatures for the alloys 41 Table 3.12. Metallography sample polishing stages 49 Table 4.1. Flow lengths for different silicon levels, superheats and section 52 thicknesses Table 4.2. Increase in flow length for increase in silicon levels, superheats 53 and section thickness Table 4.3. Metal velocity for different alloys at each superheat 53 Table 4.4. Increase in mean PSC size for increase in silicon levels, superheats 63 and section thickness Table 4.5. Increase in maximum PSC size for increase in silicon levels, superheats 63 and section thickness

x Table 4.6. Height of the silicon depleted zone in the runner 65 Table 4.7. PSC size and counts at different levels in the runner 65 Table 4.8. Effect of CMA on flow length for 6.25 mm thick section 68 Table 4.9. Effect of CMA on flow length for 12.50 mm thick section 69 Table 4.10. Height of PSC depleted zone for Al-22%Si alloy at different superheat 71 Table 4.11. Effect of phosphorus on the flow length on 6.25 mm thickness 72 Table 4.12. Effect of phosphorus on the flow length on 12.50 mm thickness 72 Table 4.13. Effect of phosphorus on the mean size of PSC measured at the gate 74 Table 4.14. Effect of phosphorus on the maximum size of PSC measured at the gate 74 Table 4.15. Height of PSC depleted zone for Al-22%Si alloy at different superheat 75 Table 4.16. Results for micro hardness tests on different phases 76 Table 4.17. Surface hardness of flow tip for different alloys and superheat 77 Table 4.18. Surface hardness at gate for different alloys and superheat 77 Table 4.19. Effect of CMA on the surface hardness at flow tip for 12.50 mm section 79 Table 4.20. Effect of phosphorus on the surface hardness at gate for 12.50 mm section 79 Table 4.21. Identification of the precipitating phases 83

xi LIST OF FIGURES

Description Page# Figure 2.1. The Al-Si binary phase diagram showing approximate 4 liquidus temperatures Figure 2.2. Feeding and flow mechanism of liquid metal 9 Figure 2.3. Flow stoppage of a pure metal in a fluidity channel 12 a) Liquid metal enters the channel and solidification begins immediately b) Solid continues to grow as metal flow c) Flow stoppage occurs as a result of ‘Choking’ at the channel entrance Figure 2.4. Flow Stoppage for hypoeutectic and eutectic alloys 13 a) Liquid enters flow channel and fine grains nucleate at tip b) Nucleation continues and fine grains grow rapidly as flow progresses c) Flow ceases when a critical concentration of solid is reached near the tip Figure 2.5. Flow stoppage mechanism in hypereutectic Al-Si alloys 14 Figure 2.6. The flow tip as observed after flow stoppage 14 Figure 2.7. Cooling curve (top) and 1st derivative curve (bottom) for A319 Al-Si alloy 16 Figure 2.8. Effect of cooling rate on the average PSC size of Al-23%Si alloy 26 Figure 2.9. Different signals from an electron beam 27 Figure 3.1. Research plan 28 Figure 3.2. A CAD sketch of the flow test pattern 34 Figure 3.3. Side view of the runner of the test pattern 34 Figure 3.4. Front view of the test pattern on the 12.50 mm attached side 35 Figure 3.5. Mixer for mixing the coating slurry 37 Figure 3.6. Positioning of the thermocouples on the runners and spirals 38 Figure 3.7. “Speedymelt” the natural gas fired furnace used for melting metal 42 Figure 3.8. The Rotary argon degasser in use 43

xii Figure 3.9. AlScan analyzer used to measure the hydrogen content in the melt 44 Figure 3.10. Data acquisition system 46 Figure 3.11. The UMSA setup (University of Windsor) 47 Figure 3.12. UMSA test sample setup (University of Windsor) 47 Figure 3.13. Buehler’s omnimet enterprise image analysis system 49 Figure 4.1. Effect of silicon and superheat on flow length 52 Figure 4.2. Effect of silicon and superheat on metal velocity 54 Figure 4.3. Schematic diagram of a casting of Al-14%Si alloy poured at 756ºC 55 showing the flow tip and the gate area Figure 4.4. Partly filled flow length (Al-22%Si alloy poured at 820ºC) 55 Figure 4.5. Effect of superheat and silicon level on undulated flow length 56 Figure 4.6. Micrographs of flow tips of alloys 57 (a) Al-14%Si (b) Al-18%Si (c) Al-22%Si Figure 4.7. Mean size of PSC at flow tip 58 Figure 4.8. Microstructure at gate area for the alloys 58 (a) Al-14%Si (b) Al-18%Si (c) Al-22%Si Figure 4.9. Eutectic silicon morphology 59 (a) Needle-like (b) Partially modified (0.7 wt % CMA) lamellar Figure 4.10. PSC morphology 60 (a) Al-14%Si shows polyhedral shaped PSC (b) Al-18%Si shows the PSC are big and the shape changed from polyhedral to star shape and eyelets start appearing (c) Al-22%Si shows the PSC are indeed big in size and in star, dendrite and polyhedral shapes Figure 4.11. Dendritic growth of PSC from wall 61 (a) left wall and (b) right wall

xiii Figure 4.12. Mean particle size for 6.25mm(left) and 12.50mm (right) 62 Figure 4.13. Max particle size for 6.25mm(left) and 12.50mm (right) 62 Figure 4.14. Change in PSC shape and size from 64 (a) bottom (0mm) to (h) top (31.75mm) in the runner Figure 4.15. Maximum melt temperature along the flow length of 12.50 mm 66 thick section Figure 4.16. Flow stoppage mechanism with interlocked PSC 67 (a) Al-14%Si, at 130 mm from gate (b) Al-18%Si, at 190 mm from gate (c) Al-22%Si, at 250 mm from gate Figure 4.17. Comparison of flow length for unmodified and CMA modified alloys 69 Figure 4.18. CMA modification in Al-14%Si alloys 70 (a) Unmodified needle-shaped eutectic silicon (b) Lamellar eutectic silicon modified with 0.7% wt CMA (c) Finely distributed eutectic silicon modified with 1.4% wt CMA (d) Fibrous eutectic silicon modified with 2.1% wt CMA Figure 4.19. Dendrites at flow tip of CMA modified Al-18%Si 71 (a) Primary and tertiary dendrites, (b) An unmodified PSC along with dendrites Figure 4.20. Comparison of flow length for unmodified and phosphorus 73 modified alloys Figure 4.21. Comparison of mean and maximum PSC size for unmodified and 74 phosphorus modified alloys Figure 4.22. PSC modification by phosphorus in Al-22%Si alloy 75 (a) Unmodified (b) Phosphorus modified Figure 4.23. Hardness values at the flow tip and gate area 78 Figure 4.24. Thermal curves for Al-14%Si (a) Cooling curve with the first derivative curve for Al-14%Si 80 with identification of precipitating phases (b) Cooling curve with the solid fraction curve for Al-14%Si 81

xiv Figure 4.25. Thermal curves for Al-18%Si (a) Cooling curve with the first derivative curve for Al-18%Si 81 with identification of precipitating phases (b) Cooling curve with the solid fraction curve for Al-18%Si 82 Figure 4.26. Thermal curves for Al-22%Si (a) Cooling curve with the first derivative curve for Al-22%Si with 82 identification of precipitating phases (b) Cooling curve with the solid fraction curve for Al-22%Si 83 Figure 4.27. Porosity in unmodified Al-18%Si alloy (a) LOM picture at 100X, (b) SEM picture at 300X 85 Figure 4.28. Identification of the phases by point EDX at the areas of interest 86 Figure 4.29. XEDS analysis of unmodified Al-22%Si alloy 87 (a) Primary silicon, gray phase 96% purity, (b) Aluminum matrix, dark phase 93% purity Figure 4.30. SEM picture of unmodified Al-18%Si alloy 87 (a) at 200X, (b) at 450X

Identification of phases by point EDX, (A) Al15(FeMn)3Si2 as white phases

appears as Chinese scripts, (B) Al2Cu also as white phases Figure 4.31. SEM picture of CMA (0.7 wt %) modified Al-18%Si alloy 88 (a) at 200X, (b) at 450X

Identification of phases, (A) Al8FeMg3Si6, as light gray phase, (B) Al2Cu as

white phases, (C) Al15(FeMn)3Si2 as white phases appears as Chinese scripts,

(D) Al5Mg8Cu2Si6 as gray phases Figure 4.32. SEM picture of CMA modified (2.1%wt) Al-18%Si 89 (a) at 200X, (b) at 450X Identification of the Mg phases with SEM-EDX

(A) Mg2Si, as black network,

(B) Al5Mg8Si6Cu2, as gray phase,

xv NOMENCLATURE

Symbol Description Units .

a Pattern thickness cm A Surface area sq. cm C Chvorinov’s constant dimensionless

CL Heat capacity of liquid metal J/g ˚C dT/dt First derivative of temperature with respect to time ºC/s fs R % solid fraction at rigidity dimensionless fs % solid fraction dimensionless fsC % solid fraction at coherency dimensionless fsR % solid fraction at rigidity point dimensionless

HE Decomposition energy of foam pattern J/g 2 hg Heat transfer coefficient W/cm ˚C

HL Latent heat of molten metal J/g

Lf Flow length cm n Mold constant in Chvorinov’s equation dimensionless t Total solidification time s

Tend End of solidification temperature ºC

Ti Ambient temperature ˚C

Tliq Liquidus temperature ºC

TM Melting temperature of metal ˚C

TR Rigidity point temperature ºC

Tsol Solidus temperature ºC v Velocity of liquid metal cm/s V Volume of Casting cu.cm ΔT Superheat ˚C

3 ρl Density of liquid metal g/cm

3 ρp Density of foam pattern g/cm

xvi ABBREVATIONS

Phrase Description . Al-14%Si Aluminum 14 wt % Silicon Al-18%Si Aluminum 18 wt % Silicon Al-22%Si Aluminum 22 wt % Silicon Al-Si Aluminum-Silicon Al-Si-Cu Aluminum-Silicon-Copper CFH Cubic Feet per Hour cm Centimeters mm Milimeters CMA Ceramic Master Alloy EDS Energy Dispersive Spectrometry EPS Expendable Poly Styrene ESI Electron Spectroscopic Imaging GFN Grain Fineness Number HRB Rockwell Hardness in “B” Scale HRC Rockwell Hardness in “C” Scale LFC Lost Foam Casting ppm Parts per million PSC Primary Silicon Crystals s Seconds SE Secondary Electrons SEM Scanning Electron Microscope UBC University of British Columbia Unmod Unmodified Alloy of the given silicon level UTS Ultimate Tensile Strength XEDS X-Ray Energy Dispersive Spectrometry YS Yield Strength α –Al Primary α –Aluminum

β-phase β – Al5FeSi Phase

xvii DEFINITIONS

Term Definition . Flow length Total portion of the spiral fully or partly filled with metal measured from flow tip to the gate

Filled flow length Fully filled (sharp edged) portion of the flow length

Partly filled Incompletely filled portion of the casting with hills and valleys flow length

Báume Unit for measurement of viscosity of water based washes at room temperature.

First derivative First derivative at each point of the cooling curve is equal to the curve slope of the cooling curve.

Solid fraction Fraction of solid (%fs) may be defined as the percentage of solid phase(s) that precipitates between liquidus and solidus temperatures in a solidifying melt.

Solid fraction Fraction of solid (%fs) when the flow of liquid metal ceases by itself at rigidity due to decrease in fluidity

xviii Chapter – 1

INTRODUCTION

Aluminum alloys play an important role in the automotive and aerospace industries, mainly due to their low density, good corrosion resistance and excellent workability. There is an ongoing interest in replacing the steel components with aluminum alloys in automotive and aerospace industries in order to reduce the overall vehicle weight. Such efforts enable improved fuel economy. In particular, there is a renewed interest in hypereutectic Al-Si alloys. These alloys show excellent wear resistance, which make them suitable for engine cylinder heads and other high wear components [1]. Suitable alloying with copper and magnesium can further improve their mechanical properties and wear resistance. In such alloys, the wear resistance is a result of finely precipitated primary silicon crystals (PSC), which are harder than the aluminum matrix. The major problem in casting hypereutectic Al-Si alloys is the control of the morphology of the PSC. Conventional casting methods such as sand casting promote segregation and excessive growth of PSC due to slow cooling. Lost foam casting (LFC) is a relatively new process and is endothermic in nature. LFC of hypereutectic Al-Si alloys has not been investigated in detail. In the LFC process [2], the foam pattern degrades, becomes liquid and then evaporates. The foam pattern is replaced by the solidifying casting, which is a replica of the foam pattern in all detail. LFC enables casting multi-layered castings as one unit, thus eliminating welding of layers. In other words, the LFC process can be used to cast intricate or complex shapes at significantly low cost. However, the endothermic nature of LFC has problems. For example, it is difficult to cast thin sections by the LFC process. In particular, hypereutectic Al-Si alloys pose unique problems. It is necessary to enable good fluidity through control of the silicon level and superheat of the liquid melt. A better understanding of the behavior of the principal phases (e.g., primary silicon and eutectic silicon) can enhance fluidity and fillability of these alloys. This can

1 result in suitable modification of the morphologies of these phases, again with a view to obtaining complete castings.

RESEARCH OBJECTIVES

The objectives of this research are stated below. i) To study the effect of silicon and superheat on alloy fluidity: Alloys with three silicon levels (14%, 18% and 22%) were studied at three levels of superheat (60ºC, 115ºC and 170ºC). The morphology of unmodified eutectic silicon and PSC were studied under these conditions. ii) To study the flow stoppage mechanism: The maximum temperature profile and image analysis was done along the metal flow length. This enabled location of the interlocked PSC zone for the alloys, which contributed to the flow stoppage. iii) To determine of the effect of modifiers on eutectic silicon and PSC morphology: Effectiveness of two silicon modifiers, ceramic master alloy (CMA) and phosphorus in modifying the morphology (shape, size and distribution) of the eutectic silicon, as well as the PSC, was determined. The precipitating phases (due to addition of modifiers) were identified.

2 Chapter – 2

LITERATURE REVIEW

In this chapter, the basic Al-Si alloys, and in particular, the hypereutectic alloys are described. The LFC process is explained. The main parameters of fluidity, such as flow length and metal velocity are then explained. Principles of feeding and solidification in metal casting are discussed with focus on feeding mechanism, solidification and castability. Flow stoppage mechanism is explained for pure metals, hypoeutectic alloys and hypereutectic alloys. Thermal analysis of solidifying alloys is described through a discussion on cooling curve analysis and Chvorinov’s rule. Various phases formed during solidification of typical Al-Si alloys are then explained. Primary and eutectic silicon phases are discussed with particular focus on modification and consequent effects on mechanical properties. The tools for characterization, in particular SEM and XEDS are discussed. Finally, the reviewed literature is related to the objectives of this thesis.

2.1. ALUMINUM ALLOYS Aluminum-copper alloy with 7% – 8% copper was the first aluminum alloy to be cast and used in aerospace industries. However, due to poor castability (mainly high shrinkage porosity), the focus shifted towards aluminum-silicon alloys, which possess better fluidity. The most widely used aluminum casting alloys are the aluminum-silicon alloys. Other aluminum-based alloys have also been developed, namely aluminum-copper-silicon, aluminum- magnesium, aluminum-zinc-magnesium and aluminum-tin alloys. All these alloys, however, invariably contain elements like manganese, magnesium, titanium and iron as alloying elements or as impurities. Alloying elements and impurities to some extent affect castability and mechanical properties of the alloy. Binary phase diagrams consisting of two major alloying elements (e.g. Al-Si alloy) are the most commonly referred to. Ternary phase diagrams are also available for alloys with three major alloying elements (e.g. Al-Si-Cu alloy).

3 2.1.1. ALUMINUM-SILICON PHASE DIAGRAM A phase diagram allows us to understand the basics of the solidification and phase transformations of an alloy. It can suggest which phase solidifies first and what phases are expected to follow. It also helps to read the solidus (Tsol) and the liquidus (Tliq) temperatures, for a given composition. However, these temperatures will vary subject to cooling rate and alloying elements present in the alloy. The Al-Si binary phase diagram is given in Figure 2.1 below.

Figure 2.1. The Al-Si binary phase diagram showing approximate liquidus temperatures [3]

From the phase diagram, the liquidus temperature for the alloys Al-14%Si, Al-18%Si and Al-22%Si were calculated at the source website [4]. These temperatures were used as the liquidus temperatures for the Al-Si alloys used in the research.

2.1.2. ALUMINUM-SILICON CASTING ALLOYS Aluminum-silicon alloys are the most popular casting alloys. They may contain Cu, Mg, Fe, and Mn as other alloying elements. Copper is added, to improve strength and hardness [5],

4 while magnesium additions provide higher strength and make them heat treatable. Silicon improves fluidity of the Al-Si alloys. It is observed that the higher the silicon content the better is the fluidity but with some limitations [6, 7]. Iron reduces the fluidity [8] in aluminum-based alloys. High-silicon alloys with more than 10% Si have low thermal expansion, which makes them suitable for high-temperature applications. When silicon content exceeds 12.6%, primary silicon crystals (PSC) precipitate, which provide excellent wear resistance in these hypereutectic Al-Si alloys. Hypereutectic Al-Si alloys are gaining impetus in the automotive engine industry due to their high wear resistance. They are a potential replacement for the wear resistant cast iron sleeves in the cylinder heads, resulting in weight reduction. However, the control of shape, size and distribution of primary silicon crystals (PSC) are the major problems encountered in casting hypereutectic Al-Si alloys. Another problem encountered with casting these alloys is the floatation and segregation of the PSC in the upper portion of the casting. It is common to have copper as another major alloying element in these alloys.

2.2. LOST FOAM CASTING (LFC) PROCESS Lost foam casting (LFC) was patented by Shroyer, in 1958, however, it was not until 1964 when Flemings performed the first casting with lost foam. The LFC process gained popularity only after it was used by General Motors in 1993. It is similar to investment casting and is very popular in producing near net shape castings. Exhaustible patterns made of EPS (Expendable Poly Styrene) foams are used as investment in contrast to wax. The EPS foam pattern used is an exact replica of the required casting. The foam pattern is embedded in a compacted sand bed. On pouring liquid metal, the foam volatilizes and the metal takes the place of the foam pattern, thus producing an exact replica of the pattern. This method has several advantages. These include: Ease in casting intricate shape, which otherwise would require lots of machining Elimination of cores and parting lines thus reducing lots of defects High flexibility in pattern making since foam is easy to manipulate for different shapes

5 High dimensional accuracy is obtained, as there is no wear and tear Cost effective due to elimination of machining of the product

Some disadvantages include: Entrapment of carbon inclusions Gas porosity due to entrapment of the pyrolysis products Surface defects like elephant skin

The LFC process involves the following steps: Preparation of the EPS foam pattern Coating of the pattern Drying of the pattern Embedding the pattern in sand/mold media Vibration of the setup for sand compaction Pouring of the liquid metal into the pattern Shaking off of the final casting after solidification

2.3. FLUIDITY STUDY Many researchers [6, 9 – 13] have worked on the fluidity of alloys and have defined fluidity in various ways. In general, fluidity is the ability of the metal to flow and fill the pattern as well as allow interdendritic transfer of metal [14]. Flow length and metal velocity of an alloy can be considered to be a measure of fluidity.

2.3.1. FLOW LENGTH Flow length in a test pattern is the length that the metal could flow before stopping by itself. Flow length depends upon many factors of which the most important factors are casting process and mold media [15]. Other factors are latent heat of fusion, surface tension, viscosity, density, freezing range and superheat. Flemings' [9] defined fluidity as,

(2.1)

6 Where,

Lf = Flow length, cm 3 ρl = density of liquid metal, g/cm

HL = latent heat of molten metal, J/g

TM = melting temperature of metal, ˚C

Ti = ambient temperature, ˚C 2 hg = heat transfer coefficient, W/cm ˚C

CL = heat capacity of liquid metal, J/g ˚C ΔT = superheat, ˚C v = velocity of liquid metal, cm/s a = pattern thickness, cm However, LFC is an endothermic process and involves extraction of heat from liquid metal. Thus, the foam barrier significantly affects fluidity. Pan and Liao [11] determined the heat of foam ablation and then modified the equation as follows,

(2.2)

Where,

HE = decomposition energy of foam pattern, J/g 3 ρp = density of foam pattern, g/cm

Flow length is also affected by superheat and is given by the equation:

Lf = (ΔT + 160) / 5.2 (2.3) Where, ΔT = superheat and

Lf = flow length

A similar correlation can also be found between total solidification time and flow length as.

Lf = 0.029 t + 33.3 (2.4) Where,

7 Lf = flow length and t = total solidification time Some other factors affecting the flow length of a metal are given as follows: Silicon level: Pan and Hu [6] and later Kamble and Ravindran [7] observed the fluidity of Al-Si alloys to increase with increase in silicon level up to 18% silicon. However, Andrews and Seneviratne [16] observed that for the same superheat, alloys with 23% silicon have higher fluidities than 20% silicon and 17% silicon. Superheat: Fluidity increases with an increase in superheat [10, 14]. This is due to increased heat content in the metal. For alloys, fluidity increases with the delay in the precipitation of the first solid phase. Hence, fluidity should increase with an increase in the superheat. Section thickness: Fluidity increases with an increase in section thickness [18, 19], due to increase in volume/area ratio. Chvorinov’s rule is applicable in the LFC process [15].

2.3.2. METAL VELOCITY Fluidity can also be described in terms of metal velocity. From the equation (2.1), it is apparent that metal velocity has a direct impact on the flow length. Hence, the flow length should increase with an increase in metal velocity, other factors remaining the same. It now becomes important to understand the process of mold feeding and the solidification process during the mold filling to have a better grasp of the metal fluidity.

2.4. MOLD FEEDING AND SOLIDIFICATION The mold feeding process for any given alloy is always the same. Even though the flow stoppage mechanism differs for different alloys, the solidification process is the same and is usually a mixture of dendritic solidification and equiaxed grain formation.

2.4.1. FEEDING MECHANISM Campbell [19, 20] described five distinct mechanisms of melt feeding during solidification of an alloy with a long freezing range. The mechanisms are discussed as follows:

8 Liquid feeding: This is the stage where the liquid metal is flowing into the fluidity channel. Mass feeding: During feeding, the liquid metal starts losing heat and solid particles start forming and the melt is actually a mixture of solid and liquid. This is also termed as mushy zone and the feeding is called mass feeding. Interdendritic feeding: As solidification progresses, dendrites form and then start becoming rigid. At this stage the melt has to go through the interdendritic passages and hence called interdendritic feeding. Burst feeding: Due to pressure of liquid metal from the back, the liquid metal may burst out of the semisolid tip and flow little further. This is called burst feeding. Solid feeding: Diffusion of alloying elements in the semi-solid or solid state is called solid feeding. The mechanisms are depicted in the Figure 2.2.

Figure 2.2. Feeding and flow mechanism of liquid metal [20]

9 All melts, even water, solidify in the form of dendrites of columnar or equiaxed shape [21]. Hence, it is very important to have knowledge of dendritic solidification, in order to study the solidification process of an alloy. Some important facts to note are as follows: Shrinkage prior to dendrite nucleation is merely linear contraction of the liquid. Mass feeding continues till dendrites impinge upon each other to form solid phase network. Gravity feeding continues inter-dendritically till the dendrites ripen and path becomes smaller in cross section and more tortuous. The practical limit for interdendritic feeding under gravity is probably the solid fraction

close to that identified as the solid fraction at rigidity point, fsR. To continue feeding at solid fraction beyond the rigidity point depends on developing sufficient stress to improve capillary action or to move or collapse a network already formed. This stress can be developed internally or externally. While solidifying and contracting, the volume deficit can generate negative stresses sufficient to increase capillary forces and even collapse adjacent dendritic network, thus feeding beyond rigidity point. This is internally applied force. Externally applied metallostatic force, for instance in pressure die-casting, squeeze casting and vibrations, help in feeding beyond rigidity point.

2.4.2. CASTABILITY AND SOLIDIFICATION CHARACTERISTICS OF Al-Si ALLOYS Castability is a broad term encompassing several aspects of casting like, sample design, design of running system, melt fluidity, cooling rate, pouring rate, dissolved gases, superheat and more. Once the liquid metal is poured, it starts losing its heat and may start solidifying immediately. Good castability of an alloy is the characteristic of the alloy to allow complete filling of the mold and also allow passage of liquid metal during solidification. On the other hand a difficult casting alloy may not allow complete filling of the mold and even after filling, shrinkage defects such as hot tear and porosity are formed within the casting. Some factors that differentiate a good casting alloy from a poor casting alloy are given in the Table 2.1.

10 Table 2.1. Stages of solidification and their effects [22]

Good casting Difficult casting Differentiating Factor alloy (A390) alloy (518) First reaches “major eutectic arrest”, Reaches major Reaches rigidity OR “solid fraction rigidity point, fsR” eutectic arrest first point, fsR first Volume fraction of isothermal solidification Quite high Very low Gap between rigidity point temperature, T and R Broader range Narrow range end of solidification temperature Tend (TR – Tend) Growth of dendrites after reaching rigidity point Growth without Growth with stress and effect of thermal contraction stress stress leading to defects Solid fractions between coherency and rigidity < 40% > 40% (fsR – fsC) (= 40% for majority of alloys)

The standard compositions of Al-Si alloys A390 and 518 are given in the table 2.2.

Table 2.2. Composition of typical A390 and 518 alloys [22]

ELEMENT Alloy A390 Alloy 518

Si 16 – 18 <0.35

Mg 0.45 – 0.65 7.50 – 8 .50

Cu 4.0 – 5.0 <0.25

Fe Max 0.50 <1.80

Mn Max 0.10 <0.35

Ti Max 0.20

Zn Max 0.10 <0.15

Others Max 0.15

2.5. FLOW STOPPAGE MECHANISM IN METALS Fluidity and flow stoppage of metals and alloys significantly differ from each other depending on their purity and alloy compositions [9, 23].

11 2.5.1. FLOW STOPPAGE MECHANISM FOR PURE METALS When metal enters the channel of a fluidity test mold, solidification begins at the channel wall and continues by the growth of columnar grains with a planar interface as metal flows through the channel. Flow ceases when the columnar grains join each other at the channel entrance as shown in the Figure 2.3.

Figure 2.3. Flow stoppage of a pure metal in a fluidity channel [24] a) Liquid metal enters the channel and solidification begins immediately b) Solid continues to grow as metal flow c) Flow stoppage occurs as a result of ‘Choking’ at the channel entrance

2.5.2. FLOW STOPPAGE MECHANISM FOR HYPOEUTECTIC ALLOYS A well developed Al dendrite skeleton grows from the mold walls to the center of the spiral channel at both the entrance and at the mid-section of the spiral casting, indicating progressive solidification. The progressive solidification mode keeps the flow channel open at the center and allows the melt to flow with less resistance. Metal flow stoppage occurs at the tip (end of the casting), when the degree of sluggishness becomes high enough to cause flow stoppage as shown in the Figure 2.4. The fluidity in hypoeutectic and eutectic composition is

12 mainly governed by the degree of melt sluggishness at the flow tip, which increases with increasing solid fractions.

Figure 2.4. Flow stoppage for hypoeutectic and eutectic alloys [24] a) Liquid enters flow channel and fine grains nucleate at tip b) Nucleation continues and fine grains grow rapidly as flow progresses c) Flow ceases when a critical concentration of solid is reached near the tip

2.5.3. FLOW STOPPAGE MECHANISM FOR HYPEREUTECTIC AL-SI ALLOYS Solidification in hypereutectic Al-Si alloys start with the precipitation of PSC and the flow stoppage occurs due to interlocking of PSC. This gives rise to three distinct zones [17, 25- 27]: Zone I, the tip is silicon free zone followed by zone II consisting of densely packed fine silicon and finally zone III consisting of larger silicon particles just behind. The solidification mode in hypereutectic Al-Si alloys was discussed first by Flemings [24]. According to him, as the metal flows into the spiral channel, aluminum dendrites start forming on the sidewalls, while the silicon keeps flowing along with the melt as silicon crystals. As the metal flows, the metal front becomes rich with the PSC and at one point they start locking with each other causing cessation of the flow.

13 Even after stoppage of flow the melt from behind which is still liquid, bursts out through the PSC and comes to the front and this portion is low in silicon and is known as silicon depleted zone. This burst of metal is aided by the high latent heat evolved during the solidification of silicon and the metallostatic pressure. This stoppage of flow and reflow of metal can happen again but finally the flow stops due to cessation of liquid melt supply due to an increase in solid fraction and loss of metallostatic pressure. Some PSC are always carried away by this melt and end up at the flow tip, hence the flow tip always contains some PSC [28, 29] as shown in Figure 2.5 below.

Figure 2.5. Flow stoppage mechanism in hypereutectic Al-Si alloys

It was observed by several researchers [6, 13, 30] that small size silicon crystals were responsible for flow stoppage unlike the large non-interlocked silicon crystals, since the flow can occur around the large crystals [7, 31]. Hence the addition of primary silicon refiners resulted in a decrease of flow length. A schematic diagram of the flow tip as studied by Pan and Hu [11] is shown in the following Figure 2.6.

Figure 2.6. The flow tip as observed after flow stoppage [11]

14 Pan and Hu [11] found that at the time when the flow ceased, the temperature near the spiral entrance read 612ºC, well above the liquidus temperature of approximately 585ºC, indicating that the metal was in a completely liquid state when the flow stopped. The temperature of the three thermocouples located directly behind the flow tip were 566ºC, 578ºC and 570ºC corresponding to the solid fractions of 34%, 12% and 28%. By means of extrapolation method, the temperature of the flow tip was determined to be around 563.5ºC at instant of flow stoppage, corresponding to a solid fraction of around 41%.

2.6. THERMAL ANALYSIS The heat evolved during solidification consists of the specific heat of the liquid metal followed by the latent heat of solidification and finally the specific heat of the solid. However, the area of interest for a metallurgist lies in and around the latent heat evolution when all phase transformations take place. By definition, thermal analysis refers to a variety of techniques in which physical property (e.g. phase changes, reactions, solid fractions) of a sample is continuously measured as a function of temperature. The curve obtained as a direct plot of temperature vs. time results is the cooling curve, which can be further analyzed as required.

2.6.1 COOLING CURVE ANALYSIS A cooling curve represents the difference between the heat loss and the heat gain due to evolution of latent heat. This curve can only show the major reactions occurring during the solidification process. In order to obtain more detailed information, the first derivative dT/dt of the cooling curve is calculated, as shown in Figure 2.7. The derivative curve is closely related to the formation of solid phase during each moment of the solidification process. Positive values of the derivative indicate high rates of formation of the solid phase, while negative values correspond to low rates of solid formation [32, 33]. At the end of solidification, no further crystallization occurs, and the cooling rate decreases gradually as the solid cools to the temperature of its surroundings. The cooling curves obtained make it possible to calculate the solid fraction at any given time. It is also possible to locate and identify phase precipitations reactions.

Figure 2.7. Cooling curve (top) and 1st derivative curve (bottom) for A319 Al-Si alloy [22]

2.6.2. SOLIDIFICATION TIME AND RANGE Fluidity of an alloy largely depends on the solidification time and the solidification range of the alloy which are described as follows: Solidification time: This is defined as the time required to reach 100% solid fraction (end of solidification) f rom the start of cooling of the melt. The start time can also be taken as the time of pouring of metal and the end time as the end of solidification temperature. However local solidification time can be calculated by the widely accepted Chvorinov’s rule [34] and is given by the following equation:

(2.5) where, t = solidification time, seconds V = volume of casting, A = surface area of the casting that contacts the mold,

16 n = constant, and C = mold constant, which depends on the properties of the metal and mold and their initial temperatures. The constant n is usually 2 [35]. Chvorinov’s rule is satisfied in the LFC process [15].

Solidification range: Solidification range is described as the temperature difference between the pouring temperature and the end of solidification or 100% solid. Thus it takes into account the superheat in the melt. This factor largely affects the metal fluidity, which represents the excess time it can flow before the start of solidification.

2.7. PRECIPITATION REACTIONS DURING SOLIDIFICATION OF Al-Si ALLOYS Intermetallic compounds are chemical compounds with a fixed atomic ratio. These are always found in a fixed composition and form as a combination of two or more elements. The intermetallics formed in an alloy are same for a given composition and change with an addition of more alloying elements. A redistribution of the precipitating phases is possible with heat treatment. Two categories of Al-Si alloys are discussed here to gain knowledge of the types of intermetallic formed and the temperature and reactions involved.

2.7.1. HYPOEUTECTIC A319 ALLOY The summary of the reactions observed during solidification process for a hypoeutectic Al-Si alloy A319 (basic composition in appendix A.3.1.) is given in the Table 2.3 below.

Table 2.3. Summary of reaction during solidification of A319 alloy [22]

Stage Summary of reactions Temperature fs (%) Start of solidification and formation of -Al 604˚C (1119˚F) 0 I α Dendritic coherency point 600˚C (1112˚F) 12 Precipitation of iron intermetallic phase Liquid = α-Al + Al (FeMn) Si II 15 3 2 579˚C (1074˚F) 35 or

Liquid = α-Al + Al15(FeMn)3Si2 + Al5FeSi. Start of main eutectic reaction 49 561˚C (1042˚F) III Liquid = Al + Si + Al FeSi 5 561˚C (1042˚F) Rigidity point 51 Precipitation of complex eutectic IV 507˚C (945˚F) 94 Liquid = Al + Si + Al2Cu + Al5Mg8Cu2Si6 V End of solidification 495˚C (923˚F) 100

For this alloy the rigidity point is reached just at the start of the main eutectic rea ction, preceded by the precipitation of one (or both) of the iron containing phases Al15(FeMn)3Si2 or

Al5FeSi. These β – Al5FeSi compounds are generally present as acicular platelets or as dendrite network. These compounds are considered very critical as they reduce the ductility and fracture toughness of the alloy. The amount and size of the β – Al5FeSi depends upon the iron content and solidification rate. Al2Cu on the other hand is a eutectic phase and is commonly found as dendrite structures. The Al2Cu phase can also grow in continuation over the β – Al5FeSi network. These two phases can be separated by solution heat treatment of the alloy [36 – 38].

2.7.2. HYPEREUTECTIC A390 ALLOY The summary of reactions observed during solidification of a hypereutectic Al-Si alloy A390 (typical composition is given in table A.3.2. in the appendix) is given in the following Table 2.4.

Table 2.4. Summary of reaction during solidification of A390 alloy [22]

Stage Summary of reactions Temperature fs (%) I Start of solidification with precipitation of PSC 654˚C (1209˚F) 0 Development of α-Al dendrites 628˚C (1162˚F) 4 II Dendritic coherency point 602˚C (1116˚F) 7 Start of eutectic reaction 565˚C (1049˚F) 10

III Liquid = Al + Si + Al5FeSi Rigidity point 561˚C (1042˚F) 36

Precipitation of Mg2Si IV 540˚C (1004˚F) 75 Liquid = Al + Si + Mg2Si Precipitation of Al2Cu V 500˚C (932˚F) 89 Liquid + Mg2Si = Al + Si + Al2Cu + Al5Mg8Cu2Si6 Precipitation of complex eutectic VI 499˚C (930˚F) 97 Liquid = Al + Si + Al2Cu + Al5Mg8Cu2Si6 VII End of solidification 494˚C (921˚F) 100

It can be observed that two types of silicon precipitate during solidification, the prima ry silicon and the eutectic silicon. In hypoeutectic alloy eutectic silicon precipitates first and in hypereutectic alloy the primary silicon precipitates first.

2.8. SILICON PHASES During solidification of hypereutectic Al-Si alloys the first phase to precipitate is the primary silicon, which precipitates as crystals. The shapes and sizes of these PSC are quite unpredictable and are dependant on silicon level of the alloy and the solidification time.

2.8.1. EUTECTIC AND PSC MORPHOLOGY Unmodified eutectic silicon can be found as coarse plates, needle like structures called acicular silicon. They are present in unpredictable [39] orientations and sizes. Unmodified primary silicon crystals (PSC) can be observed in several shapes as follows:

19 star shaped, identified as having five axes originating from a single nuclei, several variation can be found in this shape, polyhedral shaped, which appears to be hexagonal or octahedral, dendritic shaped, which resembles the dendritic form The shape and size of the PSC depend on factors like freezing rate [40], temperature gradient and local liquid composition. All these factors are dynamic during solidification and hence the shape, size and distribution of the PSC are unpredictable.

2.8.2. FLOATATION OF PSC During solidification process of hypereutectic Al-Si alloys primary silicon is the first phase to precipitate. Due to lower density of the primary silicon (specific gravity 2.33) with respect to the liquid metal (specific gravity 2.74), the primary silicon starts floating and this phenomenon is quite significant in hypereutectic Al-Si alloys [41, 42] especially with silicon more than 20%. This segregation of primary silicon depends upon the solidification time and presence of primary silicon modifier. The longer the solidification time, greater is the segregation of primary silicon and lesser is the effectiveness of the silicon modifier. Also, the segregation will increase with an increase in the silicon level of the alloy. It was observed that increasing the silicon level beyond the 16% to 19% range results in widening of the solidification range. Such widening of the solidification range would normally be expected to increase the floatation and contribute to non-uniformity of primary silicon. When casting large components, such as engine blocks, floatation of unmodified primary silicon into the risers of sand castings results in a non-uniform distribution of primary silicon and therefore detracts from the wear resistance of the alloy, which is not favored. On the other hand, modification of the PSC could lead to reduced floatation.

2.9. SILICON MODIFICATION Unmodified Al-Si hypereutectic alloys have inconsistent shape, size and distribution of the PSC as well as the eutectic silicon. Modification of the eutectic silicon and PSC can result in better distribution and consistency in shape and size and hence, may result in improved mechanical properties.

20 Silicon modifiers are added directly to the melt after degassing. However it is important to ensure that: The melt is held at the minimum possible temperature The melt is clean of scavenging impurities The melt is fluxed after addition of the modifier The melt is stirred and held for homogenization The amount of modifier to be added has to be properly decided as less addition may lead to under-modification and excess addition leads to over-modification. Over-modification of a melt can result in: rejection of the modifier in front of solidifying melt and formation of an intermetallic compound formation of coarser silicon structure and reversion of fibrous silicon into interconnected plate form Modifier fading: Though the action of the modifier may not be lost with a few remelts but the effect diminishes subsequently [43]. This phenomenon is called modifier fading. Some modifiers (like antimony) do not undergo fading; however, most of them do so with subsequent remelting. Sodium is found to be most susceptible to fading due to its high reactivity. Incubation period: This is the holding time after the addition of the modifiers for the effect of the modifier to take place. It is observed that the effect of the modifier increased with an increase in the holding time. Typically ten minutes is standard incubation period for most of the alloys. Upon modification, the microstructure of the eutectic silicon changes from coarse plate like acicular silicon to fine fibrous or lamellar silicon. However, it is not possible to achieve one hundred percent modification but effective modification will show large regions of modification. The modification level depends upon five factors, namely, type of modifier used impurities (e.g. phosphorus) present in the melt amount of modifier used freezing rate silicon content of the alloy

21 2.9.1. EUTECTIC SILICON MODIFICATION MECHANISM Eutectic silicon is present in needle-like shapes, which are detrimental to the mechanical properties of the alloy. The eutectic silicon modifiers are capable of modifying the shape of the eutectic silicon from needle-like to lamellar and further to fine fibrous structure [32, 44]. However, excess addition of modifier may lead to formation of plate-like structures as a result of over-modification.

2.9.2. PRIMARY SILICON MODIFICATION MECHANISM With the unmodified primary silicon, the size and distribution of the PSC are very inconsistent with chances of getting PSC depleted zones and PSC rich zones. Modification of these PSC can result in better distribution and consistency in shape and size of the PSC. The modifiers provide numerous additional sites for nucleation of the primary silicon thus leading to more number of PSC and hence smaller size PSC. This also leads to a better distribution of PSC provided the metal was homogeneous. The amount of modifier to be added has to be properly determined as less addition may lead to under-modification and excessive addition may lead to rejection of phosphorus, which is then carried along the flow and remains concentrated at the flow tip.

2.9.3. EFFECTS ON MECHANICAL PROPERTIES The mechanical properties of aluminum-silicon alloys can be improved by inducing structural modifications of the normally occurring eutectic silicon. Addition of certain elements, such as calcium, sodium, strontium, and antimony, to the hypo and hypereutectic aluminum- silicon alloys results in a finer lamellar or fibrous eutectic network. Typically, modified structures exhibit higher tensile properties and appreciably improved ductility [45] when compared to similar but unmodified structures. A study by Hengcheng et al. [46] with Al-Si alloy and strontium addition concluded that: 1. With an addition of 0.0375% Sr to Al–11.6% Si alloy, there was a significant increase in the amount of dendritic α -Al phase by 236.9% compared to the unmodified alloy. 2. With an increase of Sr content from 0 to 0.0375%, there was increase of ultimate tensile strength and fracture elongation by 34.0 and 714.6%, respectively. Also that, a combination measure of mechanical properties increased by 100%.

22 3. It is also found that, the mechanical properties of the modified alloy is linearly related to the amount of dendritic α-Al phase, and that the α-Al phase plays an important role in improving the mechanical properties of near-eutectic Al–Si alloys. A study by Pennors et al. [47] have indicated that the addition of phosphorus leads to an increase in the amount of β-needles in the structure with the AlP particles were found to nucleate along the sides of the needles. However, superheating the melt at 900˚C for 1 hour dissolves a large portion of the AlP and hence lesser amount of β-phase formation. Addition of Sr thereafter, negates the beneficial effect of superheat treatment, with AlP particles again appearing with an increase in β-phase.

2.9.4. TYPES OF SILICON MODIFIERS There are two types of silicon present in hypereutectic Al-Si alloys, the primary silicon and the eutectic silicon. The primary silicon is almost pure silicon and is present as crystals of different shapes and sizes. The eutectic silicon is the needle shaped acicular silicon. Modifi cation of both the silicon may result in highly improved mechanical properties. Several silicon modifiers are commonly used in the industry however, it is important to note here that some are effective in modifying the eutectic silicon and others are effective on primary silicon. The most common eutectic silicon modifiers are elements of groups IA, IIA and some rare earth elements. Of these, sodium is most effective and very widely used followed by strontium and antimony. These modifiers are added at very low concentration levels of 0.01% to 0.02%. Rapid solidification too can result in a modified structure. Phosphorus is an impurity in modification of hypoeutectic Al-Si alloys since it reacts with the modifiers especially Na. However in hypereutectic Al-Si alloys it is deliberately added in hypereutectic alloys as silicon modifier. The phosphorus reacts with Al to form AlP, which act as nucleates for the primary silicon.

2.9.4.1. CERAMIC MASTER ALLOY (CMA) CMA, patented by UBC, is a composite of nano-sized alumina particles dispersed in Mg or Al-Mg alloy. These nano-sized alumina particles act as hetero-nucleation sites during alloy solidification in the alloy to be modified. The mechanism of grain refinement is believed to be the heterogeneous nucleation sites offered by the nano-sized alumina particles.

23 CMA was manufactured and characterized at the University of British Columbia and its role as a suitable grain refiner and modifier was assessed at the University of Windsor (UW) and industry partners (NEMAC). Trials done at UBC on A319 cast aluminum alloy showed that CMA (added 0.7% by weight) could refine grains and modify the eutectic silicon. Comparisons were m ade with Sr (as modifier) and Ti-B, Ti-C (as grain refiner). As compared to these two, CMA had better thermodynamic stability in the alloy, and did not appear to fade. Alumina, being a stable oxide, did not dissolve in the melt and thus the grain refining capability did not fade away with reuse. This effect was studied [48] by multiple remelting and casting of the CMA modified A319 alloy. Preliminary data indicated that there was a substantial improvement in the mechanical properties of the cast aluminum alloy as a result of modification. The addition of CMA resulted in: A microstructure with grain refinement up to 60%, A decrease in the SDAS A partial modification of the eutectic Al-Si phase An increase of 25% in the Rockwell hardness of CMA grain refined A319 A significant improvement in the YS and UTS of the aluminum alloy. (There was a 48.5% and 76.1% increase in the average YS and UTS values at a strain rate of 0.001s-1 and 18.2% and 36.5% improvement at 0.0001 67 s-1 respectively.)

Addition of magnesium [49 – 52] resulted in two precipitations, Mg2Si which

precipitated as black dots along the side of eutectic silicon and transformation of the β – Al5FeSi

into Al8Mg3FeSi6. Furthermore, these additions led to an increased amount of β – Al5FeSi needles in the structure, where the AlP provides heterogeneous nucleation sites [53] for silicon and was observed to nucleate along the sides of the β – needles. This phenomenon was overcome by dissolution of the AlP, by superheating and holding the melt at 900ºC for one hour. However, addition of Sr negated the effect of superheat treatment [47].

2.9.4.2. PHOSPHORUS Unmodified primary silicon remains as chunks randomly distributed in the matrix along with the eutectic silicon. These big silicon chunks impart an inhomogeneous hardness over the surface. To achieve homogeneous hardness distribution, it is important to break the big silicon chunks into smaller size and evenly distribute within the matrix [54].

24 Phosphorus was first used and patented by Sterner Rainer [55]. Later, Onitsch-Modl [56]

and Mascr´e [57] observed refining of primary silicon, by addition of PCl5 and Cu-P alloy to an Al alloy containing between 17 and 23-wt% Si. It has been pre-established that phosphorus [58] interferes with the eutectic silicon modification mechanism. Phosphorus reacts with sodium and probably with strontium and calcium to form phosphides (Na3P) thus nullifying the intended modification additions. Hence, it is desirable to use low-phosphorus metal when modification is a process objective, also make an excess amount of modifier additions to compensate for phosphorus-related losses. Addition of phosphorus more than 0.001% (10 ppm) resulted in coarsening of eutectic structure and plate- like primary silicon phase. While phosphorus less than 0.0003% – 0.0005% (3 – 5 ppm) showed a refined eutectic phase and a finely divided lamellar structure [59 – 61]. AlP also acts as nucleation sites for iron bearing phases. According to Loper, Jr. and Cho [41], AlP that forms, promotes formation of blocky eutectic silicon particles, inhibiting formation of silicon flakes. However in hypereutectic alloy phosphorus additions improve the silicon distribution and improve machinability. Xiangfa et al. [62] observed that addition of boron along with phosphorus yielded better silicon distribution and machinability than phosphorus alone. Table 2.5 shows the effect of phosphorus modification of the PSC into smaller size.

Table 2.5. PSC size in Al-22%Si alloy treated with 0.055%P [63]

Casting Measured PSC size (μm)

Conventional pressure die casting 20

3/16 inch plate poured in copper mold 15

1 inch diameter bar in cast iron mold 29

1 ½ thick block in cast iron mold 48

1 ½ thick block in sand cast 78

25 The PSC size was also affected by the cooling rate. The effect of cooling rate combined with the addition of phosphorus resulted in further refinement of the PSC, which is shown in the Figure 2.8.

Figure 2.8. Effect of cooling rate on the average PSC size of Al-23%Si alloy [64]

2.10. SCANNING ELECTRON MICROSCOPY (SEM) AND X-RAY ENERGY DISPERSIVE SPECTROMETRY (XEDS) The scanning electron microscope (SEM) is a microscope that uses electrons rather than light to form an image. It is possible to study many features using SEM. These include topography, morphology, composition and crystallography of the samples. Topography is a study of surface features of an object and the texture. The detectable features are limited to a few nanometers. Morphology is the study of shape, size and arrangement of the particles in the specimen, lying on the surface of the sample or has been exposed by grinding or chemical etching. The detectable features are limited to a few manometers. XEDS helps finding the composition of elements and compounds the sample is composed of and their relative ratios, in areas ~ 1 micrometer in diameter. The crystallographic information gives the arrangement of atoms in the specimen and their degree of order.

26 Chapter – 3

EXPERIMENTAL PROCEDURE

This chapter provides a research plan followed by a detailed description of materials, pattern making, experiment setup, and the melting and pouring procedures for producing the castings. Further, it describes the procedures for analysis of castings, which includes the sampling procedures and the testing methodology.

3.1. RESEARCH PLAN The aim of this research was to study the fluidity, flow stoppage and silicon modification in three hypereutectic aluminum silicon alloys with 14%, 18% and 22% silicon. The fluidity was studied for three superheats of 60ºC, 115ºC and 170ºC, with the addition of two silicon modifiers ceramic master alloy (CMA, patented by UBC) and phosphorus. The flow stoppage was studied by examining the casting under light optical microscope to locate the interlocked PSC region. Modification of the eutectic silicon and PSC were studied with respect to their morphology. Floatation of PSC and the effect of modifiers on the same were also studied. Analysis of the castings was done by studying the flow length, metal velocity, eutectic silicon morphology, PSC morphology, cooling curves, first derivative curves, phase identification using XEDS and finally studying the micro and surface hardness of the castings. A detailed research plan is given in the Figure 3.1.

27 RESEARCH PLAN

A Study of Fluidity, Flow Stoppage and Silicon Modification in Lost Foam Casting of Hypereutectic Al-Si Alloys

Al-14%Si Alloy Al-18%Si Alloy Al-22%Si Alloy

60°C Superheat 115°C Superheat 170°C Superheat

Phosphorus CMA CMA (2.1 wt%) + (0.015 wt %) Modification Phosphorus (0.015 wt Modification %) Modification

0.7 wt % 1.4 wt % 2.1 wt %

Analysis of Castings

Fluidity Image Thermal SEM/ Hardness XEDS

Flow Metal Cooling First Micro Surface Length Velocity Curve Derivative Hardness Hardness

Eutectic PSC SEM Phase

Silicon Mophology Imaging Identification

Morphology

Figure 3.1. Research plan

28 3.2. MATERIALS The consumables used in this research were Al-Si alloys, silicon modifiers, EPS foam boards and pattern coating.

3.2.1. Al-Si ALLOYS The base alloy used for the experiment was A319 hypoeutectic aluminum silicon alloy. An aluminum-silicon master alloy (Al-50%Si) was mixed with A319 alloy to achieve the targeted silicon levels of 14%, 18% and 22%. The compositions of the alloys, A319 and Al-50%Si are given in the Tables 3.1 and 3.2, respectively.

Table 3.1. Composition of A319 alloy used for the experiment [Arnberg et al. 1996]

Composition (wt %)

ELEMENT STANDARD ACTUAL

Si 5.5 - 6.5 6.15

Mg Max 0.1 0.096

Cu 3 - 4 3.65

Fe Max 1.0 0.47

Mn Max 0.5 0.31

Ti Max 0.25 0.14

Zn Max 1.0 0.71

Others Max 0.15 0.50

29 Table 3.2. Composition of the Al-50%Si hardener alloy

Composition (wt %)

ELEMENT ACTUAL

Si 53

Fe 0.17

Ti 0.02

V <0.01

B <0.01

Other Elements <0.05

3.2.2. SILICON MODIFIERS Two types of silicon modifiers were used in the research, ceramic master alloy (CMA, patented by UBC) and copper-phosphorus brazing alloy. CMA is a composite of nanosize alumina particles suspended in magnesium or aluminum-magnesium alloy. The CMA addition was done in three levels 0.7 wt %, 1.4 wt % and 2.1 wt %. The composition and general properties are given below (Tables 3.3. and 3.4.). Small chunks of CMA were cut using a hacksaw and used as modifier additions to the melt.

Table 3.3. Ceramic master alloy (CMA) composition

Components Amount (wt %) Magnesium Or 60 – 93 Alloy of magnesium (70%) and aluminum (30%)

Aluminum oxide powder 7 – 40

30 Table 3.4. Ceramic master alloy (CMA) general information

Properties Description

Physical state Ingot

Color Silver

Magnesium 650ºC Melting point/ range Magnesium-aluminum alloy 470ºC

Boiling point/ range 1107ºC

Density: at 20ºC 1.8 – 3.0 g/cm3

Phosphorus was added (0.015 wt %) in the form of cut pieces of copper-phosphorus brazing wire. It was added at 0.015 wt %. The composition of the Cu-P brazing alloy is given below (Table 3.5).

Table 3.5. Fos Flo-7, copper-phosphorus brazing alloy composition

Composition (wt %)

ELEMENT % ACTUAL PRESENT

Cu 92.75

P 7.25

3.2.3. EPS FOAM BOARDS The patterns were prepared from 1.6 pcf foam boards. With a view to avoiding inhomogenity, an entire spiral was prepared from one board. The EPS foam board specifications are given below (Table 3.6).

Table 3.6. Specifications of the EPS foam

Property Specifications

Foam grade Nova F271T

Foam density 1.6 pcf

Additives None

3.2.4. COATING Mica base low permeability high absorption coating was used to coat the pattern (Table 3.7).

Table 3.7. Specifications of the coating

Coating parameters

Low Permeability Quality High Absorption

Coating Type Mica-base

Median particle size (μm) 30 – 50

Percentage of Solids 43.80 – 44.80%

Brookfield Viscosity at 20 RPM (cP) 1100 – 1300

Báume 40

Thixotropic Index 1.80

GM Perm Air Flow(cm3/(cm2-s)) 1.0-6.0

SRI Saturation capacity (%) 0.4 – 0.5

SRI Saturation Rate (g/(in3-s)) 5.0 – 10.0

SRI Dry Density (ρ) 0.75 – 0.85

32 3.2.5. MOLD MEDIA Low density unbonded synthetic mullite (AFS GFN # 35) was used as the mold medium. Before each use the mullite quality was checked for excessive burnt material and for agglomerated particles. After every two experiments, approximately half the volume of used mullite was replaced by fresh unused material.

3.3. PATTERN MAKING The spiral pattern making consisted of three steps. These consisted of cutting and shaping of the spirals and runner, coating of the spirals and runner and assembly of the final test pattern. The EPS foam pattern consisted of two spirals attached to a common runner and a hollow fiber sprue. The two spirals, runner and the sprue were attached with the help of a craft glue gun. This is a fillet type joint, and thus there was no glue on the joining surface. This joint ensured minimum loss of heat and velocity at the joints. The dimensions of the spirals [2, 6, 31] and the runner are given in Table 3.8. The dimensions were tested to satisfy Chvorinov’s rule [34].

Table 3.8. Dimensions of the spirals

Part Dimensions

Total length 1080.00 mm (42.50 inch) each

Width 12.50 mm (0.50 inch) Spirals

Thickness 12.50 mm (0.50 inch) and 6.25 mm (0.25 inch)

Runner 127.00 mm X 50.00 mm X 31.75 mm

For preparing the spirals, a template was designed using Solidworks. Using a CNC machine, the spiral template was then cut from a polypropylene sheet. The spirals were cut using this template.

The pattern design used for the research is given in the Figure 3.2.

Figure 3.2. A CAD sketch of the test pattern

The spirals were attached to the runner as shown in the following Figures 3.3 and 3.4. This ensured that the bases of the spirals were at the same height of 10.00 mm from the bottom and side ends of the runner.

Figure 3.3. Side view of the runner of the test pattern showing the position of the spirals

Figure 3.4. Position of the 12.50 mm spiral on the runner

To prepare the spiral test pattern, first, a sketch of the spiral was drawn using the template. Four spirals were drawn on a board using only the center portion of the board [65] thus ensuring uniform surface finish and bead fusion. A rough spiral shape was then cut using a hot wire cutter. These rough-cut spirals were again cut into specific thickness of 6.25 mm and 12.50 mm with the hot wire cutter. To ensure uniformity, the required sections were cut

35 from either surfaces of the 31.75 mm thick foam board, while the middle portion was discarded. These thin section rough-cut spirals were then attached to the spiral template using an adhesive tape. The spirals were given a finishing cut with a high-speed router to ensure smooth surface. A runner of given dimensions (Table 3.8) was cut from the same EPS foam board. The spirals and runner were then coated with the mica base low permeability high absorption coating and dried (Section 3.2.1). The dried runners and spirals were joined using a glue gun. It was ensured that there was no coating left on the joining surfaces. Excess coating was removed using a knife. The coated spirals were then attached to the coated runner. Finally a hollow fiber sprue, 115.00 mm high was attached at the center of the runner (Figure 3.4).

3.3.1 COATING AND DRYING The individual pattern parts were coated with the mica base low permeability high absorption coating. The coating was mixed using a coating mixer (Figure 3.5) and the viscosity was measured. A viscosity corresponding to a Báume number of 40 was obtained. The slurry was then transferred into a larger vessel and the pattern components were dip coated with a single dip. The coated pattern components were then left hanging for the excess coating to drip off and ensure uniform coating thickness [66, 67]. This process was reinforced by rotating the runner at a very low speed. The pattern was left in the open air for 4 to 6 hours and then dried in an oven at 54.4ºC (130 °F) for 24 hours.

Figure 3.5. Mixer for mixing the coating

3.4. EXPERIMENT SETUP The experimental setup involved three steps. These were preparation and position of thermocouples, embedding of the pattern assembly in sand and compaction, and connecting the thermocouples to the data acquisition system.

3.4.1. PREPARATION AND POSITION OF THERMOCOUPLES Sixteen ‘K’ – type thermocouples were prepared in the lab, using a 12.50 mm long two-hole insulating sleeve (diameter 3.18 mm). The wires were passed through the holes of

37 the insulators and fused at the open end with help of a flame to form a bead 0.4 – 1.0 mm in diameter. The positions of the thermocouples on the pattern assembly were as follows: Two thermocouples were placed on either side of sprue on the centerline of the runner at a distance of 16.25 mm from the edge (Figure 3.6).

Figure 3.6. Positioning of the thermocouples on the runners and spirals

In addition to the two above-mentioned thermocouples, fourteen thermocouples were placed along the outer side of both the patterns at the following distances from the gate as given in Tables 3.8 and 3.9 for the two section thicknesses of 12.50 mm and 6.25 mm, respectively.

38 Table 3.9. (a). Position of thermocouples on 12.50 mm thick spiral

Thermocouple Distance from gate # 3 25.4 mm # 4 88.9 mm # 5 152.4 mm # 6 215.9 mm # 7 279.4 mm # 8 342.9 mm # 9 406.4 mm # 10 444.5 mm

Table 3.9. (b). Position of thermocouples on 6.25 mm thick spiral

Thermocouple Distance from gate # 11 25.4 mm # 12 88.9 mm # 13 152.4 mm # 14 215.9 mm # 15 279.4 mm # 16 342.9 mm

To place a thermocouple, a small hole of the same diameter as the thermocouple insulator and about 3.5 mm deep was drilled on the side of the spiral manually at specified locations. The thermocouple was then pushed all the way down till the end of the hole. Particular care during inserting the thermocouples ensured that they did not come out of the pattern later.

39 3.4.2. PATTERN EMBEDDING AND SAND COMPACTION The casting setup was done in a 508 mm (20 inches) diameter flask on the vibration table. The dried pattern assembly was placed in the vibrating flask containing synthetic mullite. The pattern assembly (spirals, runner and sprue) was leveled horizontally. It was also ensured that the pattern was placed at the same depth from the top of flask every time to ensure same level of compaction [68]. Sixteen ‘K type’ thermocouples were inserted. The flask was then very slowly filled with mullite sand and, as the level reached about one inch above the pattern assembly, the flask was vibrated horizontally at an acceleration of 0.7G for 20 seconds. The flask was again filled with mullite to the height of the pouring cup. The flask was then vibrated horizontally at 1.0G for 40 seconds. A pouring cup was then fixed on the top. This two-stage vibration was chosen so as to ensure that the thin section patterns do not break and the thermocouples do not come out of the spirals.

3.4.3. CONNECTION TO DATA ACQUISITION Each of the thermocouples was then connected to the data acquisition system with care, to ensure proper matching of the thermocouple number and the position. The default readings were checked for errors with a view to corrections to loose connection or wrong polarity. The data acquisition system collected data in sequence from all the points and worked at a frequency of 8 Hz.

3.5. MELTING AND ALLOY PREPARATION Alloys of composition Al-14%Si, Al-18%Si and Al-22%Si were prepared with fresh ingots of A319 alloy and Al-50%Si master alloy. The weight of the A319 ingots and the required Al-50%Si alloy were calculated using MS-Excel software. The compositions of the alloys obtained are given in the Appendix A.3.1 – A.3.3. Typical sequence of activities during melting and pouring operation is given in the Table 3.10.

Table 3.10. Sequence of activities and temperature schedules

Activity Observation/Temperature

Double preheat for 10 minutes if starting from cold Add smaller pieces of metal and start furnace, keep

adding the rest metal regularly Add flux Metal appears molten

Stir the melt for homogenization After it appears molten All materials are molten or Turn furnace “OFF” target temperature is reached Skim slag Pour temp + 70ºC

Degas using rotary argon degasser for 210 seconds Pour temp + 70ºC

Analyze hydrogen using AlScan

Add silicon modifier, stir, add flux and hold

Take out crucible from furnace Pour temp + 10ºC

Start data acquisition system

Start pouring Pour temp [Table 3.3]

Three alloys were poured at three different superheats and the details are given in Table 3.11.

Table 3.11. Pouring temperatures for the alloys

Alloys Pouring temperatures

Al-14% Si 655ºC 710ºC 765ºC

Al-18% Si 710ºC 765ºC 820ºC

Al-22% Si 765ºC 820ºC 875ºC

Superheat 60ºC 115ºC 170ºC

All the materials were melted in a gas-fired furnace “Speedymelt” (Figure 3.7).

Figure 3.7. “Speedymelt” the natural gas fired furnace used for melting alloys

3.5.1. DEGASSING AND MEASUREMENT OF HYDROGEN CONTENT Before degassing, the melt surface was cleaned with a clean scoop coated with boron nitride. As the required temperature of the melt was reached, degassing was done with the help of the rotary argon degasser. This degasser was built at the laboratory by the author and a colleague N. Dahata with constructive participation from other research group members. Precautions were taken before using the degasser. First, the lance was coated with boron nitride and preheated with an electric heat gun at 200ºC for at least 1 hour before every degassing operation. The lance was rotated at a very slow speed to achieve uniform heating. The argon supply was then attached to the argon inlet of the degasser and the argon cylinder supply valve was opened. As the melt was ready, the degasser was then taken out of the stand and placed on the furnace lid so that the lance was at the center of the crucible. Initially, the lance was rotated at a very slow speed and the argon flow increased using the control knob on the argon regulator, so that small bubbles evolved at the surface. The lance speed was then

42 increased and adjusted with the help of the variable DC voltage controller, till the melt stirring was just below turbulent. Care was taken that impurity from metal top was not pulled into the melt. Further adjustment of the argon flow was done so that evenly distributed fine bubbles appeared at the melt surface. Degassing was done for 210 seconds with a lance speed between 300 – 400 rpm and argon flow between 7 and 10 CFH (Figure 3.8). After degassing, the hydrogen content of the melt was checked using the AlScan hydrogen analyzer (Figure 3.9). The AlScan was calibrated as per the manufacturers’ manual. According to the quality requirements outlined by the “Near-net-shape processing of materials” laboratory, if the hydrogen content was above the targeted range of 100 – 150 ppm, the melt had to be degassed again. However, such a necessity did not arise, and the standard procedure (outlined above) resulted in the target value of hydrogen in the melt. The hydrogen in the degassed melt was in agreement with the data of other researchers [69].

Figure 3.8. The rotary argon degasser in use

Figure 3.9. AlScan hydrogen analyzer used to measure the hydrogen content in the melt

3.5.2. ADDITION OF MODIFIERS After achieving the required hydrogen content and the melt was at the required temperature (Table 3.10), the silicon modifiers were added. CMA was added as chunks and phosphorus was added in the form of cut Cu-P brazing wires. The melt was then stirred gently, fluxed and held for about seven minutes before pouring.

3.6. POURING PROCEDURE After completion of the melt treatments and analysis, the crucible was taken out of the furnace. The melt temperature was checked again before pouring into the mold. The metal was then slowly poured into the pouring cup and filled up to about 30 cm from the top. The speed was adjusted to maintain the same melt height to ensure the same metallostatic head pressure during pouring. All the pours were repeated twice and if

44 necessary they were repeated for a third time. This ensured that the results obtained were consistent and free of experimental errors.

3.7. ANALYSIS OF CASTINGS Analysis of the castings was done using thermal analysis, macro analysis and micro structural analysis. The flow tip, gate and the region of the flow spiral with anticipated primary silicon blockage were specifically investigated. Chemical analysis was carried out using the optical emission spectrometer (OES) at “Gamma Foundry” in Markham, Ontario.

3.7.1. MEASUREMENT OF FLUIDITY Fluidity was studied from the points of flow length and metal velocity. The flow length for this experiment is defined as the length of the spiral filled with metal, referred to as ‘cast spiral’, taking the gate (joint between spiral and runner) as point ‘zero’. The length of the cast spiral was measured using a measuring tape for each section thickness. The flow length given is an average of three readings. Metal velocity was calculated using the thermocouples. Figure 3.6 shows the exact location of the thermocouples in the spirals. The velocity was calculated using equation 3.1 below.

V = 63.5 / (T2 – T1) (3.1) where,

T1 = Metal arrival time at the first thermocouple (at 25.4 mm from gate)

T2 = Metal arrival time at the second thermocouple (at 63.5 mm from first thermocouple)

3.7.2. THERMAL DATA ANALYSIS The thermal data accumulated by the data acquisition system (Figure 3.10) was analyzed [70] using Matlab and MS-excel. The cooling curve data, first derivatives and fraction solids were analyzed using the “universal metallurgical simulator and analyzer” (UMSA) at the University of Windsor.

Figure 3.10. Data acquisition system

3.7.3. UMSA The “universal metallurgical simulator and analyzer” (UMSA), was used to arrive at the cooling curves, the first derivative curves and solid fractions [33, 71, 72] for a slow steady state cooling rate. The data acquisition rate was set at 5 Hz and the UMSA thermal analysis software was used to derive the cooling curves and the first derivative curves. The runs were repeated for the second time for each alloy. The UMSA setup with the testing chamber is given in the Figure 3.11.

Figure 3.11. UMSA setup [University of Windsor]

The test chamber has the option to treat the sample in vacuum and also with options for variable cooling rates. The heating is done by an induction coil while the cooling can be under vacuum, in open air, with air jet or water quench. The sample heating setup for the sample is given in the figure 3.12.

Figure 3.12. UMSA test sample setup [University of Windsor]

47 3.7.4. SAMPLING FOR ANALYSIS Sampling was done in such a manner that all tests were done on the same sample so that the results could be correlated. First, the flow tips were cut at appropriate lengths to examine the partly filled flow lengths. The top surfaces of these samples were then ground to get a smooth surface and the macro hardness tests carried out. The base was also ground to make it flat. The samples were then cut to 50 mm length from the tip and further used for image analysis after preparation, as outlined in section 3.6.6. The spiral was cut off from the runner at the gate and 50 mm long samples were cut off from the gate side for hardness tests and image analysis similar to the flow tip samples. Later, these samples were used for SEM analysis and finally for the micro hardness tests. For locating the interlocking PSC zone, a sample about 100 mm long was cut from the identified area and investigated under the light optical microscope.

3.7.5. HARDNESS TESTING Macro and micro hardness tests were carried out on the flow tips and at the gate area. The surface hardness tests were carried out on the tips and the gate on the flow spiral, as well as the runner. The samples for hardness tests were ground at the top surface. This ensured that the flow tip was flat and had an even surface; the gate area grinding was matched up for uniformity in the sampling procedure. The macro hardness tests were carried out using the Buehler automatic macro hardness tester. Micro hardness of the PSC and aluminum matrix was tested on samples from the gate area and was done using the Buehler micro hardness tester. The samples for these tests required good polishing and hence were done on the samples prepared for image analysis.

3.7.6. SAMPLE PREPARATION FOR IMAGE ANALYSIS Samples for image analysis were prepared after the macro surface hardness tests were carried out. Small and thin samples were cold mounted. The metallographic samples were prepared using manual methods and the steps adopted are elaborated in the Table 3.12.

48 Table 3.12. Metallographic sample polishing stages

Step Material Media Lubricant Remarks #

I Grinding paper # 120 Grit Water Flat and smooth surface

II Grinding paper # 320 Grit Water Old scratches removed

III Grinding paper # 600 Grit Water Shiny appearance

IV Polishing paper # 1200 Grit Water Mirror like surface Diamond paste Diamond No scratches observed under V Polishing cloth # 6 micron paste extender microscope Diamond paste Diamond VI Polishing cloth # 1 micron paste extender

Optical micro analysis was don e with digital imaging using Buehler’s “Omnimet enter prise image analysis system” (Figu re 3.13).

Figure 3.13. Buehler’s Omnimet enterprise image analysis system

49 3.7.7. MEASUREMENT OF PSC SIZE Using the optical image analyzer with Omnimet software, the mean and maximum size of the PSC was measured. The Omnimet software used first measured the area covered by the particle, for a magnification of 100X and they appear as highlighted portions. The software then calculated the equivalent spherical diameter of the highlighted portion. This diameter was expressed as the size of the particle in this study. Mean particle size was calculated as an average diameter of all the calculated diameters. The final result was an average of twenty readings and readings were taken at ten different locations.

3.7.8. SEM ANALYSIS After image analysis, the polished samples were subjected SEM analysis. The aim of SEM analysis was to identify the precipitating phases and correlate with the first derivative (dT/dt) curve of the alloy. Phase identification procedure: The identification of phases was done using SEM-EDX. The identification process involved the following two steps, A) X-Ray mapping: Here the area of interest is scanned and the element rich portions are identified, so it becomes clear as to what phase to expect. B) Point XEDS: This is done at the point where the phase has to be identified. For

example, to locate the Al2Cu, the portion or point, which showed rich in copper and aluminum, was examined with point EDX. This point EDX performs a spectrometric analysis at that point and calculates the percentage compositions by weight and atomic ratio.

50 Chapter – 4

RESULTS AND DISCUSSION

This chapter presents the results as they relate to fluidity, flow stoppage and silicon modification in three hypereutectic Al-Si alloys cast by the lost foam casting (LFC) process and is broadly divided into seven sections. The first section explains the significant effect of section thickness, superheat and silicon levels on flow characteristics, silicon morphology and primary silicon crystals (PSC) floatation, in the unmodified Al-Si alloys. The contribution of interlocked PSC to the flow stoppage in hypereutectic alloys is discussed in the next section. In the third section, it is demonstrated that the ceramic master alloy (CMA) is effective in modifying the eutectic silicon but does not affect the PSC, thus having no effect on metal flow length and PSC floatation. It is shown in section four that phosphorus modified the PSC, and this resulted in reduced flow length and PSC floatation. The surface hardness increased with increases in silicon level and with additions of CMA and phosphorus, as shown in section five. In the next section, the cooling curves and first derivative curves are used to illustrate the various phases precipitating during the solidification process. Finally, in section seven, the precipitated phases are identified using SEM-XEDS.

4.1. FLOW CHARACTERISTICS OF UNMODIFIED AL-SI ALLOYS The flow characteristics of the unmodified Al-Si alloys with respect to silicon levels and superheat were compared in terms of flow length, metal velocity and surface undulations at flow tips.

4.1.1. FLOW LENGTH The flow length for the Al-Si alloys at three silicon levels, subject to different section thickness and superheat were measured and tabulated in the Table 4.1 and the results depicted in Figure 4.1 for better comparison. Aluminum-14 wt% Silicon, Aluminum-18 wt% Silicon, Aluminum-22 wt% Silicon will be referred to as Al-14%Si, Al-18% Si, Al-22% Si, respectively. The flow length data in the table are an average of three readings.

51 Table 4.1. Flow lengths for different silicon levels, superheats and section thicknesses

Flow Length (mm)

Superheat (ºC) 60 115 170 Alloys Thickness (mm) 6.25 12.5 6.25 12.5 6.25 12.5

MEAN 136 209 140 269 154 342 Al-14%Si Range ± 10 ± 15 ± 10 ± 15 ± 10 ± 15

MEAN 187 348 210 374 222 397 Al-18%Si Range ± 10 ± 20 ± 10 ± 20 ± 10 ± 20

MEAN 204 361 259 400 302 446 Al-22%Si Range ± 15 ± 25 ± 15 ± 25 ± 15 ± 25

Figure 4.1. Effect of silicon and superheat on flow length

It was observed that the flow length increased with an increase in silicon level and superheat. With higher superheat, the precipitation of the first phase primary silicon in this case was delayed. Thus, the melt flowed for a longer period before the start of solidification. The increase in the flow length with increasing silicon level was due to the high latent heat of solidification of silicon (264 cal/gm). This latent heat provided enough heat to maintain fluidity of the alloy. The increase in flow length due to increase in superheat and silicon level are given in the Table 4.2.

52 Table 4.2. Increase in flow length for increase in silicon levels, superheats and section thickness

Increase in flow length (%) For superheat increase from 60ºC For silicon increase from 14% to Parameter to170ºC at each silicon level 22% at each superheat level Al-14%Si Al-18%Si Al-22%Si 60ºC 115ºC 170ºC 6.25 mm section 13 19 47 50 85 96 12.50 mm section 63 15 23 72 50 30

Examining the effect of section thickness, with an increase in silicon level, the flow length for the 6.25 mm thick section increased by 50-96%, while for the 12.50 mm, flow length increased by 30-72%, for the superheats studied. Thus, the effect of silicon level was more significant in the 6.25 mm thick section. This can be attributed to the latent heat of solidification of silicon and increase in the alloy solidification range. Thus, with a higher cooling rate, a higher silicon level provides more sensible heat to maintain fluidity.

4.1.2. METAL VELOCITY The metal velocity was calculated using the first two thermocouples on each spiral as described in section 3.7.1. The results are tabulated in the Table 4.3.

Table 4.3. Mean metal velocities and superheats

Mean metal Velocity (mm/sec)

Silicon level Al-14%Si Al-18%Si Al-22%Si

Thickness 6.25mm 12.5mm 6.25mm 12.5mm 6.25mm 12.5mm

60ºC 9.03 13.60 10.08 15.01 11.01 16.59

Superheat 115ºC 9.64 16.77 11.38 19.41 13.40 20.02

170ºC 11.25 18.93 11.92 22.26 15.16 23.03

The above values of metal velocity are depicted in Figure 4.2 for comparison.

Figure 4.2. Effect of silicon and superheat on metal velocity

It was observed that the mean metal velocity increased with an increase in superheat and silicon level. This was possibly due to a delay in the precipitation of the first phase (PSC in this case). Thus the melt was able to flow for a longer period before reaching the rigidity solid fraction. The increase in metal velocity with an increase in silicon level is attributed to the increase in the solidification range for these alloys with increase in silicon level. High latent heat of fusion of silicon is another factor affecting the solidification range.

4.1.3. STUDY OF FLOW TIP AND GATE AREA The flow tip surface and the microstructure at the flow tip and gate area for the alloys were studied. The following is a schematic representation of the flow tip and the gate area (Figure 4.3).

Figure 4.3. Schematic diagram of a casting of Al-14%Si alloy poured at 765ºC, showing the flow tip and the gate area

The total flow length consisted of fully filled (sharp edged) portion, which will be referred to as the “filled flow length”, and a partly filled (hills and valleys) flow portion, which will be referred to as “partly filled flow length” (Figure 4.4). The total flow length (filled flow length + partly filled flow length) will be referred to as “flow length”. This phenomenon was observed for all levels of silicon and superheats.

Figure 4.4. Partly filled flow length (Al-22%Si alloy poured at 820ºC)

55 4.1.3.1. Flow tips The partly filled flow length was measured and the mean values are shown below (Figure 4.5).

Figure 4.5. Effect of superheat and silicon level on partly filled flow length

It was observed that the length of partly filled flow was affected mainly by the silicon level and was not significantly affected by superheat. It is quite possible that the stoppage of flow and subsequent re-flow of metal occurred at a particular temperature for a given composition of the alloy since every alloy has a fixed temperature for solid fraction at rigidity. Hence, pouring temperature did not affect the partly filled flow length. The microstructure of the flow tip revealed that it primarily consisted of eutectic silicon in aluminum matrix. Only a few PSC were found at the flow tip that seems to have drifted along with the metal re-flow, following flow stoppage as discussed in section 2.5.3 earlier. The microstructure of flow tips for different silicon levels are shown in Figure 4.6 taken at a magnification of 50X.

Figure 4.6. Micrographs of flow tips of alloys (a) Al-14%Si, (b) Al-18%Si, (c) Al-22%Si

The PSC mean size was measured for all the alloys (Figure 4.7). The mean PSC size is an average of 20 readings taken as per the procedure given in section 3.6.7 earlier.

Figure 4.7. Mean size of PSC at flow tip

From the measurements, it was observed that the mean PSC size increased with an increase in silicon level, which is attributed to the increase in the alloy solidification range (hence more time for growth of PSC). There was also an increase (7.5% to 19%) in size due to an increase in the section thickness from 6.25 mm to 12.50 mm. This is attributed to a decreased cooling rate [73] for the 12.50 mm section thickness.

4.1.3.2. Gate Area The microstructure at the gate revealed that, unlike the flow tip, the gate area consisted mainly of PSC as shown in figure 4.8 below [50X magnification].

Figure 4.8. Microstructure at gate area for the Al-22%Si alloy

58 There was a significant change in the PSC morphology (shape, size and distribution) with an increase in silicon level. Further discussion on this is given in section 4.2.2 later.

4.2. SILICON MORPHOLOGY IN UNMODIFIED Al-Si ALLOYS The effect of superheat and silicon level was related to the shape and size of PSC and eutectic silicon. It was observed that the eutectic silicon morphology remained unaffected by the superheat and the silicon level. However, there was a significant effect on the PSC shape and size with an increase in silicon level and superheat.

4.2.1. EUTECTIC SILICON Unmodified eutectic silicon appears as needle-like structures (acicular silicon), and these are known to have unpredictable orientations and sizes [39]. The growth of eutectic silicon depends on factors such as cooling rate, modifiers and the local melt composition. These factors change continuously during solidification and hence the morphology of the eutectic silicon is unpredictable. However, upon modification, the shape of eutectic silicon can change from needle-like to lamellar to a fine fibrous structure. Figure 4.9 below shows unmodified needle-like eutectic silicon and partially modified (with 0.7 wt % CMA) lamellar eutectic silicon. The microstructures of the flow tip area are seen at a magnification of 50X using the Buehler image analyzer.

Figure 4.9. Eutectic silicon (a) Needle-like (b) Partially modified (0.7wt%CMA) lamellar

59 4.2.2. SHAPE, SIZE AND DISTRIBUTION OF PRIMARY SILICON As the silicon level increased in the alloy from 14% to 18% and 22%, the shape, size and distribution of PSC changed (Figure 4.10) [50X magnification].

Figure 4.10. PSC morphology (a) Al-14%Si shows polyhedral shaped PSC, (b) Al-18%Si shows big PSC and the shape changed from polyhedral to star shape and eyelets start appearing, (c) Al-22%Si shows the PSC are very big in size, in star, dendrite and polyhedral shapes

At 14% silicon, the PSC shape was polyhedral. As the silicon level increased to 18%, a few PSC changed to star shape and eyelets started appearing in the PSC with an increase in size. At 22% silicon, the PSC size was indeed big and the shapes were polyhedral, star and dendrite. The shape and size of the PSC depend on factors such as freezing rate, temperature gradient and local liquid composition. All these factors continuously change during solidification and hence the morphology of the PSC is unpredictable. However, at a higher

60 silicon level, the PSC sizes are indeed large. This is due to an increase in the solidification range leading to growth of the PSC. The change in the shape from polyhedral to elongated shapes can be attributed to the directional cooling affects. Figure 4.11 show the dendrite shaped PSC growing from the wall on the runner. The micrographs are at a magnification of 50X.

Figure 4.11. Dendritic growth of PSC from wall of the runner (a) left wall and (b) right wall

It was also observed that the PSC tend to grow perpendicular to the edge, in the direction opposite to the heat flow. This again shows that the PSC tend to grow dendritically and hence sensitive to direction of heat flow. The PSC size was measured for all the levels of silicon, superheat and section thickness. The mean size of PSC and the maximum size of PSC for each case are given in Figures 4.12 and 4.13. The mean PSC size is an average of 20 readings taken as per procedure given in section 3.6.7 earlier.

Figure 4.12. Mean particle size for 6.25mm (left) and 12.50mm (right) at gate area

Figure 4.13. Max particle size for 6.25mm (left) and 12.50mm (right) at gate area

It was observed that both PSC mean size and PSC maximum size increased with an increase in silicon level for both section thicknesses. The increase in mean PSC size and maximum PSC size are shown below (Table 4.4 and 4.5).

62 Table 4.4. Increase in mean PSC size for increase in silicon levels, superheats and section thickness

Increase in mean PSC size (%)

Superheat increase from 60ºC Silicon increase from 14% to 22% Parameter to170ºC

Al-14%Si Al-18%Si Al-22%Si 60ºC 115ºC 170ºC

6.25 mm section 17 15 6 99 83 80

12.50 mm section 17 11 10 94 96 82

Table 4.5. Increase in maximum PSC size for increase in silicon levels, superheats and section thickness

Increase in maximum PSC size (%)

Superheat increase from 60ºC to Silicon increase from 14% to 22% Parameter 170ºC Al-14%Si Al-18%Si Al-22%Si 60ºC 115ºC 170ºC

6.25 mm section 18 28 33 192 213 227

12.50 mm section 16 45 40 300 309 383

The effect of superheat on the particle size was not so significant for 14% silicon but significantly increased for 22% silicon. The maximum PSC size almost doubled from 14% to 18% silicon and then again from 18% to 22% silicon. This again is attributed to the increase in the solidification range with an increase in silicon level.

4.3. THE PSC FLOATATION PHENOMENON Due to the difference in density of aluminum (specific gravity 2.74) and silicon (specific gravity 2.33), silicon has a tendency to float to the top portion of the casting. This phenomenon was observed more significantly with an increase in solidification time and is known as the silicon floatation process.

63 4.3.1. PSC DISTRIBUTION ACROSS THE CROSS-SECTION HEIGHT In the test spiral pattern, PSC floatation was clearly observed on the 31.75 mm thick runner for the Al-22%Si alloy. PSC floatation was not observed on the thin (6.25 mm) section. However, on the 12.50 mm thick section, the PSC floatation seemed to be present but was not distinguishable. PSC floatation however, was not observed for Al-14%Si and Al-18%Si alloys. Figure 4.14 shows the PSC distribution along the height for Al-22%Si alloy.

Figure 4.14. Change in PSC shape and size from the (a) top (30.0 mm) to (h) bottom (0.0 mm) on the runner [50X magnification]

Though individual PSC floated to the top of the runner, some PSC adhered to the wall and appeared to grow in size. The size of the PSC also increased towards the top. The PSC

64 depleted zone expressed as the height of that zone increased with an increase in superheat. The heights of the PSC depleted zone at different superheats are given in Table 4.6.

Table 4.6. Height of the silicon depleted zone in the runner of Al-22%Si alloy

Superheat 60ºC 115ºC 170ºC

Height (mm) 4.9 5.9 7.0

Clearly, the height of the PSC depleted zone increased with an increase in superheat. This was because the PSC had more time to float upwards. Also, due to buoyancy, bigger size PSC rose faster than the smaller size particles. Thus, the degree of floatation increased with the solidification time, and increased with the silicon level of the casting.

4.3.2. SIZE AND DISTRIBUTION OF PSC The PSC size and counts at different levels on the runner of Al-22%Si alloy at different superheat levels are given below (Table 4.7), as per procedure given in 3.6.7 earlier.

Table 4.7. PSC size and counts at different levels of the runner

Level 60ºC 115ºC 170ºC

Maximum Size 25 mm – Top 508.2 560.0 711.1

of PSC 18 mm – 25 mm 257.1 301.3 373.3 (microns) 10 mm – 18 mm 112.4 119.9 138.4

25 mm – Top 83.9 87.6 93.4 Mean Size of 18 mm – 25 mm 64.4 65.1 69.4 PSC (microns) 10 mm – 18 mm 46.2 46.8 47.8

25 mm – Top 59 56 51

Counts of PSC 18 mm – 25 mm 106 94 85

10 mm – 18 mm 194 158 152

65 It was observed that the maximum size and the mean size of PSC increased from bottom to the top of the runner. Again, the number of count of PSC decreased from the bottom to the top of the runner. This is the result of the PSC floatation in the Al-22%Si alloy. It can also be observed that the maximum size of the PSC increased with an increase in superheat. Thus, the floatation of silicon is significantly dependant on the time of solidification.

4.4. EFFECT OF PSC ON FLOW STOPPAGE The maximum temperature curve of the solidifying melt was plotted relative to the distance from the gate for the different silicon levels (Figure 4.15). A plateau was observed for all the curves. A sample from this plateau region was examined under the optical microscope.

Figure 4.15. Maximum melt temperature along the flow length of 12.50 mm thick section

66 The micrographs revealed the flow stoppage mechanism with interlocked PSC blocking the flow passage, leading to the initial flow stoppage. However, the metal again flowed as the passage reopened due to evolution of heat of fusion of silicon and the metallostatic head pressure from the liquid metal behind. This evolution of heat due to very high latent heat of fusion of silicon (264 cal/gm) compared to aluminum (95 cal/gm) led to an increase in the temperature of the region resulting in the plateau on the maximum temperature curve. Figure 4.16 shows the interlocked PSC, responsible for flow stoppage.

Figure 4.16. Flow stoppage mechanism with interlocked PSC (a) Al-14%Si, at 130 mm from gate, (b) Al-18%Si, at 190 mm from gate, (c) Al-22%Si, at 250 mm from gate [50X magnification]

Furthermore, it was observed that the size of the interlocking PSC increased with an increase in the silicon level. However, the smaller interlocked PSC were more effective in

67 flow stoppage than the bigger ones. As a result, the flow after the initial flow stoppage was higher with higher silicon level. Thus, the partly filled flow length was observed to increase with an increase in the silicon level. Again, the distance of the interlocking PSC area from the gate increased with an increase in silicon level. With increased silicon level, the total flow length, the partly filled flow length and the fully filled flow length were longer.

4.5. EFFECT OF CMA ON FLUIDITY AND ALLOY MICROSTRUCTURE The effects of addition of CMA on the alloys were studied with respect to flow length and silicon modifications. CMA was added initially at 0.7 wt % and then at 1.4 wt % followed by 2.1 wt % to study the effect on silicon modification.

4.5.1. FLOW LENGTH The flow length for the unmodified alloy was determined for each section thickness, silicon level and superheat. The resulting flow lengths are shown below (Tables 4.8 and 4.9 and Figure 4.17). The values are an average of three readings.

Table 4.8. Effect of CMA on flow length for 6.25 mm thick section

Flow length (mm)

Al-14%Si Al-18%Si Al-22%Si Super- heat % % % Unmod CMA Unmod CMA Unmod CMA Change Change Change 60ºC 136 133 2.2 207 204 1.4 219 217 0.9

115ºC 140 137 2.1 210 209 0.5 259 263 -1.5

170ºC 154 155 -0.6 212 215 -1.4 302 298 1.3

68 Table 4.9. Effect of CMA on flow length for 12.50 mm thick section

Flow length (mm)

Al-14%Si Al-18%Si Al-22%Si Super- heat % % % Unmod CMA Unmod CMA Unmod CMA Change Change Change 60ºC 209 210 0.5 348 350 0.6 361 367 1.7

115ºC 269 265 -1.5 374 368 -1.6 400 392 -2.0

170ºC 342 345 0.9 397 394 -0.8 446 442 -0.9

Figure 4.17. Comparison of flow length for unmodified and CMA modified alloys

It was observed that with the addition of CMA, there was a nominal change (± 2.1% for 6.25 mm thick section and ± 2.2% for 12.50 mm thick section) in the flow length. The CMA was not effective in modifying the PSC, and hence the flow length was not affected.

4.5.2. EUTECTIC SILICON MORPHOLOGY CMA modified the eutectic silicon from needle-like to lamellar structure. With increase in the amount of CMA, the eutectic silicon changed further to a finer eutectic structure. The change in eutectic silicon structure from needle-like to lamellar to fine fibrous structure can be seen in Figure 4.18 taken at a magnification of 50X.

Figure 4.18. CMA modification in Al-14%Si alloys (a) Unmodified needle-shaped eutectic silicon, (b) Lamellar eutectic silicon modified with 0.7 wt% CMA, (c) Finely distributed eutectic silicon modified with 1.4 wt% CMA, (d) Fibrous eutectic silicon modified with 2.1 wt% CMA

With the modification of the eutectic silicon into a fine fibrous structure, the dendrites appeared more prominent. Figure 4.19 shows the dendrites in the Al-14%Si alloy modified by the CMA. The PSC were not modified even with a significant (2.1 wt%) amount of CMA addition. The sample shown was taken from the flow tip of Al-18%Si alloy at a magnification of 50X.

Figure 4.19. Dendrites at flow tip of CMA modified Al-18%Si (a) Primary and tertiary dendrites, (b) An unmodified PSC along with dendrites

4.5.3. PSC MORPHOLOGY CMA did not have any effect on shape, size or distribution of PSC.

4.5.4. PSC FLOATATION Addition of CMA in all the three different ratios did not affect the PSC floatation as shown in Table 4.10. CMA did not modify the PSC. It can be seen that for Al-22%Si alloy the PSC depleted zone was unaffected.

Table 4.10. Height of PSC depleted zone for Al-22%Si alloy at different superheat

Height measured for the alloys (mm)

Superheat Al-22%Si Al-22%Si + CMA 60ºC 4.9 4.9 115ºC 5.9 5.8 170ºC 7.0 6.9

71 4.6. EFFECT OF PHOSPHORUS ON FLUIDITY AND ALLOY MICROSTRUCTURE Phosphorus was added as copper-phosphorus brazing alloy at 0.015 by weight. The effects of phosphorus addition were studied and the results and discussions follow.

4.6.1. FLOW LENGTH The flow length values for each silicon level and superheat are given below (Tables 4.11 and 4.12).

Table 4.11. Effect of phosphorus on the flow length on 6.25 mm thickness

Flow length (mm)

Al-14%Si Al-18%Si Al-22%Si Super- heat % % % Unmod Phos Unmod Phos Unmod Phos Change Change Change 60 136 106 -22.1 207 167 -19.3 219 173 -21.0

115 140 112 -20.0 210 178 -15.2 259 208 -19.7

170 154 120 -22.1 212 182 -14.2 302 236 -21.9

Table 4.12. Effect of phosphorus on the flow length on 12.50 mm thickness

Flow length (mm)

Al-14%Si Al-18%Si Al-22%Si Super- heat % % % Unmod Phos Unmod Phos Unmod Phos Change Change Change 60 209 182 -12.9 348 315 -9.5 361 302 -16.3

115 269 235 -12.6 374 323 -13.6 400 342 -14.5

170 342 301 -12.0 397 347 -12.6 446 400 -10.3

The data are presented graphically in Figure 4.20.

Figure 4.20. Comparison of flow length for unmodified and phosphorus modified alloys

With phosphorus addition, the flow length reduced by 12% to 18% for 6.25 mm section and 12% to 15% for 12.50 mm section. This is attributed to the modification of PSC and reduction in their size as discussed in the following section 4.6.2. The reduction in flow length was more significant in 6.25 mm due to higher cooling rate resulting in further reduction of PSC size.

4.6.2. EUTECTIC SILICON MORPHOLOGY Phosphorus had no effect on the eutectic silicon.

4.6.3. PSC MORPHOLOGY With the addition of phosphorus, modification of the PSC occurred. There was a change in the mean and maximum size of the PSC and it also affected the PSC floatation. The mean and maximum sizes of PSC obtained (section 3.6.7) due to phosphorus addition are shown in the Tables 4.13 and 4.14. Samples were taken from the gate (PSC rich) area.

73 Table 4.13. Effect of phosphorus on the mean size of PSC measured at the gate (PSC rich)

Mean size of PSC (microns) Al-14%Si Al-18%Si Al-22%Si Super- % % % heat Unmod Phos Unmod Phos Unmod Phos Change Change Change 60ºC 42.0 36.9 -12.2 56.9 48.8 -14.3 81.4 68.9 -15.3 115ºC 42.1 38.6 -8.2 60.0 51.6 -13.9 85.2 68.9 -19.1 170ºC 49.2 40.7 -17.2 61.7 53.6 -13.2 89.8 71.2 -20.6

Table 4.14. Effect of phosphorus on the maximum size of PSC measured at the gate (PSC rich)

Maximum size of PSC (microns)

Al-14%Si Al-18%Si Al-22%Si Super- % % % heat Unmod Phos Unmod Phos Unmod Phos Change Change Change 60ºC 115.0 88.5 -23.1 161.5 128.7 -20.3 508.2 214.9 -57.7 115ºC 137.1 97.9 -28.6 168.1 130.2 -22.5 560.0 239.8 -57.2 170ºC 147.4 78.6 -46.7 168.9 130.1 -23.0 711.1 240.3 -66.2

The mean and maximum PSC size is shown (Figure 4.21) for unmodified and modified alloys by phosphorus for various superheats.

Figure 4.21. Comparison of mean and maximum PSC size for unmodified and phosphorus modified alloys

74 The micrographs below show the effect of phosphorus modification, which resulted in reduced PSC size compared to unmodified alloy in Al-22%Si alloy (Figure 4.22) at 50X.

Figure 4.22. PSC Modification in Al-22%Si alloy (a) Unmodified, (b) Phosphorus modified

Similar effects were observed in Al-18%Si and Al-14%Si with reduction in PSC size and more uniform distribution of the PSC.

4.6.4. PSC FLOATATION The height of the PSC depleted zone was measured for Al-22%Si alloy at different superheats (Table 4.15). The height of the PSC depleted zone for CMA addition is also given for comparison. Since PSC floatation did not take place in Al-18%Si and Al-14%Si, only the values for Al-22%Si are given.

Table 4.15. Height of PSC depleted zone for Al-22%Si alloy at different superheat

Height measured for the alloys (mm) Superheat Al-22%Si Al-22%Si + CMA Al-22%Si + Phosphorus 60ºC 4.9 4.9 1.0 115ºC 5.9 5.8 1.4 170ºC 7.0 6.9 2.3

The height of the PSC depleted zone increased with an increase in superheat. However, with the addition of phosphorus the height of silicon depleted zone reduced for all

75 the superheats. This was due to the fact that phosphorus modified the PSC into smaller size as discussed earlier, and smaller size PSC float slower than the bigger ones as discussed in section 4.3.1. Thus, the degree of floatation was significantly reduced with phosphorus addition.

4.7. SURFACE HARDNESS STUDIES

Hardness readings were taken on the samples to study the effect of eutectic and PSC morphology on the surface hardness of the alloys. The micro hardness of the phases was also determined.

4.7.1. MICRO HARDNESS TESTS Micro hardness tests were performed on the alloys to measure the hardness of the different phases and to relate their effect on the macro hardness of the material. The hardness measurements are given in the Table 4.16. For comparison of results hardness of 100% pure silicon is 1000-1200 HV [74], and hardness of 99.99% pure Al is 120-140 HV [75].

Table 4.16. Results for micro hardness tests on different phases

Micro hardness of phase HV HRC HRB Primary silicon crystals (96% purity) 727.2 61.1 Too high

Aluminum matrix (93% purity) 78.9 Too low 16.0

The Micro hardness results show that the hardness results were lower than those in the literature [74, 75]. This may be a result of porosity related defects, which is very common in foam casting of hypereutectic Al-Si alloys especially with high pouring temperature. Another reason could be the experimental errors, which could arise due to insufficient thickness of the PSC and cushioning effect by the aluminum matrix.

76 4.7.2. SURFACE HARDNESS OF THE UNMODIFIED ALLOYS Surface hardness tests were carried out at the gate and the flow tip on the top surface of the spirals, with a view to study the effect of silicon distribution in the casting. The hardness readings at the flow tip are given in Table 4.17.

Table 4.17. Surface hardness of flow tip for different alloys and superheat

Surface hardness (Rockwell “B” scale)

Al-14%Si Al-18%Si Al-22%Si Superheat Mean SD Range Mean SD Range Mean SD Range

60ºC 41.0 2.88 ± 5.3 44.7 1.82 ± 4.1 45.5 3.94 ± 7.6

115ºC 39.3 4.36 ± 7.9 49.0 3.46 ± 5.8 55.7 3.71 ± 7.2

170ºC 42.0 6.18 ± 7.7 52.0 3.62 ± 8.3 58.8 4.12 ± 7.3

It was observed that the surface hardness of the flow tip was a function of the silicon content in the alloy. Surface hardness increased by 16% for Al-18%Si and by 29% for Al- 22%Si with an increase in superheat. However, for the Al-14%Si alloy there was no significant change in the surface hardness of the alloy. The hardness readings at the gate area are given in the Table 4.18.

Table 4.18. Surface hardness at gate for different alloys and superheat

Surface hardness (Rockwell “B” scale)

Al-14%Si Al-18%Si Al-22%Si Superheat Mean SD Range Mean SD Range Mean SD Range

60ºC 67.0 2.48 ± 4.8 69.2 3.50 ± 6.0 74.5 4.40 ± 7.1

115ºC 67.2 3.48 ± 5.4 75.4 3.61 ± 4.9 81.9 4.85 ± 9.3

170ºC 64.2 3.33 ± 4.9 71.4 5.58 ± 9.3 86.9 5.21 ± 9.5

At the gate area, the surface hardness increased with an increase in silicon level. For Al-22%Si, there was a significant increase in hardness with an increase in superheat. This was

77 due to the silicon floatation process in Al-22%Si. The primary silicon floated to the top resulting in an increase in hardness at the top surface. Silicon floatation was not observed in Al-14%Si and Al-18%Si alloys. The surface hardness on flow tip and gate area at different silicon levels and superheat are graphed below (Figure 4.23).

Figure 4.23. Hardness values at the flow tip and gate area

The surface hardness at the flow tip was always lower compared to the gate. This is due to the fact that the flow tips consisted of eutectic silicon in aluminum matrix with few PSC [28].

4.7.3. EFFECTS OF MODIFIERS ON SURFACE HARDNESS Hardness test samples were cut from the flow tip (eutectic silicon rich area) and the gate (PSC rich area) and tested for both modifiers. The results indicated are an average of 25 readings taken for each sample.

4.7.3.1. EFFECTS OF CMA It was observed that, with addition of CMA, there was no significant change in the surface hardness at the gate (PSC rich) area. However, at the eutectic silicon rich flow tips the results showed an increase in surface hardness (Table 4.19).

78 Table 4.19. Effect of CMA on the surface hardness at flow tip for 12.50 mm section

Surface hardness (Rockwell ‘B’ scale) Al-14%Si Al-18%Si Al-22%Si Super- % % % Unmod CMA Unmod CMA Unmod CMA heat Change Change Change 60ºC 41.0 49.7 21.30 44.7 52.1 16.49 45.5 52.4 15.22

115ºC 39.3 51.2 30.18 49.0 55.6 13.38 55.7 62.4 11.90

170ºC 42.1 53.3 26.84 52.0 58.7 12.91 58.8 65.4 11.34

Surface hardness at the flow tip increased on addition of CMA because CMA modified the eutectic silicon from needle-like to lamellar uniformly distributed eutectic silicon (Figure 4.18). Since, the flow tip primarily consisted of eutectic silicon, the surface hardness increased 12% to 30% and the increase was more prominent for lower silicon level. There were also small increases in the surface hardness with increasing superheat due to silicon floatation. Similar results were obtained for the 6.25 mm section thickness.

4.7.3.2. EFFECTS OF PHOSPHORUS It was observed that with phosphorus addition, there was no significant change in surface hardness of the alloys at flow tips. However, a change in the surface hardness was observed at the gate as shown in Table 4.20.

Table 4.20. Effect of phosphorus on the surface hardness at gate for 12.50 mm section

Surface hardness (Rockwell ‘B’ scale)

Al-14%Si Al-18%Si Al-22%Si Super- % % % heat Unmod Phos Unmod Phos Unmod Phos Change Change Change 60ºC 67.0 75.8 13.2 69.2 77.7 12.3 74.5 82.8 11.1

115ºC 67.2 75.1 11.7 75.4 80.2 6.4 81.9 84.9 3.7

170ºC 64.2 73.2 14.1 71.4 77.5 8.6 86.9 85.8 -1.3

79 Surface hardness at the gate increased by 8% to 14% with addition of phosphorus (0.015 wt%), for Al-14%Si and Al-18%Si. This was due to refinement and better distribution of the PSC. However, with Al-22%Si the hardness increased by up to 11% but there was a minor decrease (1.3%) in hardness at the highest superheat. The effect of superheat on hardness was also significantly reduced. This indicates that phosphorus additions, resulted it reduced floatation of PSC and thus a better distribution of the PSC. A similar trend was observed for the 6.25 mm section thickness for both at the flow tip and the gate area.

4.8. THERMAL ANALYSIS The cooling curves were developed at the University of Windsor, using the universal metallurgical simulator and analyzer (UMSA). The following Figures 4.24 – 4.26, show the cooling curves, first derivative curves and solid fraction for the three alloys Al-14%Si, Al- 18%Si and Al-22%Si. The heating and cooling curves with the first derivative curves are in Appendix, Figures A.4.1. – A.4.3.

Figure 4.24. (a) Cooling curve with the first derivative curve for Al-14%Si with identification of precipitating phases

Figure 4.24. (b) Cooling curve with the solid fraction curve for Al-14%Si

Figure 4.25. (a) Cooling curve with the first derivative curve for Al-18%Si with identification of precipitating phases

Figure 4.25. (b) Cooling curve with the solid fraction curve for Al-18%Si

Figure 4.26. (a) Cooling curve with the first derivative curve for Al-22%Si with identification of precipitating phases

Figure 4.26. (b) Cooling curve with the solid fraction curve for Al-22%Si

The precipitating phases the reaction temperature and the solid fraction at that temperature are given in the Table 4.21.

Table 4.21. Identification of the precipitating phases

Alloys Al-14%Si Al-18%Si Al-22%Si Point Temp Temp Temp Precipitating phase (%) (%) (%) (deg C) fs (deg C) fs (deg C) fs Primary silicon A 612 0.0002 643 0.0002 681 0.0002 crystals B Aluminum as dendrite 604 0.5 605 6.9 622 14.7 Fe bearing C 580 6.6 591 8.5 581 24.2 intermetallics D Eutectic reaction 561 10.2 559 15.7 558 29.5

E Al2Cu 499 96.7 492 96.9 487 98.3

F Complex eutectic 478 99.0 480 99.2 471 99.6

G End of solidification 461 100.0 460 100.0 456 100.0

83 Solidification ranges derived for th e alloys, using the UMSA ex periment al setup, ar e as follows: Al-14%Si alloy = 151ºC Al-18%Si alloy = 183ºC Al-22%Si alloy = 225ºC It was observed that the eutectic reaction temperature slightly decreased with an increase in silicon level, but the solid fraction during eutectic reaction increased with increase in silicon level. The solidification range increased marginally with increase in the silicon level of the alloy. The end of solidification (100% solid fraction) temperature decreased with an increase in silicon level. Thus, with the same superheat the actual solidification time is much higher for higher silicon level. All these explain the improved fluidity with an increase in silicon level. The cooling rates for different alloys were calculated using the maximum temperature recorded by the thermocouple and the time at the end of solidification. The resulting cooling rates achieved for the different pattern thickness were as follows: I) 6.25 mm thick spiral – 0.66 [C/s] or 39.4 [C/min] II) 12.50 mm thick spiral – 0.44 [C/s] or 26.2 [C/min] III) 31.75 mm thick runner – 0.35 [C/s] or 20.8 [C/min]

4.9. PRECIPITATED PHASES AND INTERMETALLIC COMPOUNDS The scanning electron microscope (SEM) was used to study the precipitating phases for the alloys. Some phases that were not visible with the light optical microscope (LOM) could be seen using the SEM. Using the LOM, samples the phases were observed and identified using XEDS.

Figure 4.27 shows pictures of porosity observed under LOM and SEM.

Figure 4.27 Porosity in unmodified Al-18%Si alloy (a) LOM picture at 100X, (b) SEM picture at 300X

4.9.1. XRAY MAPPING AND XEDS X-Ray mapping was done after identifying an area of interest. The XEDS spectrum showed the elements present and their relative composition. The indications appeared as dots scattered over the area. The detection process was run for about 10 minutes and individual element windows appeared separately. Point XEDS was then done on the identified area. This gave an approximate composition of the phase. For example, to identify the iron bearing phases, a point XEDS was done on the area which showed high concentration of Fe in the mapping micrographs. Some phases identified are shown in the Figure 4.28.

Figure 4.28. Identification of the phases by point EDX at the areas of interest

4.9.2. IDENTIFIED PHASES The phases and intermetallic compounds precipitating during the solidification process were identified using SEM-EDX. The SEM pictures also revealed some more information on the phases present that were not visible with the light optical microscopy (LOM). Figure 4.29 shows the purity of the silicon and aluminum phases as analyzed by

XEDS. Figure 4.30 shows the identified iron bearing phases and the Al2Cu phases.

Figure 4.29. XEDS analysis of unmodified Al-22%Si alloy (a) Primary silicon, gray phase 96% purity (b) Aluminum matrix, dark phase 93% purity

Figure 4.30. SEM picture of unmodified Al-18%S i alloy (a) at 200X, (b) at 450X, Identification of phases by point EDX, (A) Al15(FeMn)3Si2 as white phases appears as Chinese scripts, (B) Al2Cu also as white phases

87 Additional phases with addition of CMA: Addition of CMA, which contains 40-70% magnesium or aluminum-magnesium, resulted in precipitation of few more Mg bearing phases. Figure 4.31 shows the additional Mg rich phases that were observed.

Figure 4.31. SEM picture of CMA (0.7 wt%) modified Al-18%Si alloy (a) at 200X, (b) at 450X Phases present (A) Al8FeMg3Si6, as light gray phase, (B) Al2Cu as white phases, (C) Al15(FeMn)3Si2 as white phases appears as Chinese scripts, (D) Al5Mg8Cu2Si6 as gray phases

Iron bearing phase Al15(FeMn)3Si2 was observed as Chinese scripts along with some magnesium bearing phases such as Al8FeMg3Si6 as light gray phases and Al5Mg8Si6Cu2 as gray phases. With further increase of CMA additions to 2.1 wt %, precipitation of Mg rich phase was seen in large measure (Figure 4.32).

Figure 4.32. SEM picture of CMA modified (2.1wt%) Al-18%Si (a) at 200X (b) at 450X, Identification of the Mg phases with SEM-EDX (A) Mg2Si, as black network, (B) Al5Mg8Si6Cu2, as gray phase, (C) Al8FeMg3Si6, as light gray phase

Magnesium bearing phase Mg2Si was observed with increase in CMA level. Mg2Si were found as a combination of Chinese scripts [75, 76, 77] and dendrite network [78, 79]. The darkness of these phases increased with an increase in Mg content in the phase.

4.10. COMBINED EFFECT OF PHOSPHORUS AND CMA The combined effect of phosphorus (0.015 wt%) and CMA (2.1 wt %) addition was studied. It was observed that, neither the PSC nor the eutectic silicon was modified.

89 Chapter – 5

CONCLUSIONS AND SUGGESTION FOR FUTURE WORK

5.1. EFFECT OF SILICON AND SUPERHEAT ON ALLOY FLUIDITY 1) Flow length and metal velocity increased with an increase in silicon level and superheat for both 6.25 mm and 12.5 mm thicknesses. This was due to an increased solidification range and delayed precipitation of the first phase, PSC in this case.

2) For the 6.25 mm thick section, the effect of silicon in increasing the flow length was more pronounced relative to that of superheat. For the 12.5 mm section thickness, both silicon and superheat had the same effects. The latent heat of fusion of silicon had more of an effect in supplying sensible heat than superheat, especially for the high cooling rate.

3) The flow length was not affected by CMA modifications of the alloys, but with phosphorus addition, the flow length decreased by 12% to 18% for the 6.25 mm section and 12% to 15% for the 12.50 mm section. This is because phosphorus was able to modify the PSC into a smaller size. CMA did not affect the PSC.

4) The eutectic reaction temperature slightly decreased with an increase in silicon level, but the solid fraction at the eutectic reaction temperature increased with an increase in the silicon level. The eutectic solidification range also increased marginally with an increase in the silicon level of the alloy. This improved the fluidity of the alloy.

5) The end of solidification (100% solid fraction) temperature decreased with an increase in silicon level. Thus, with the same superheat, the actual solidification time is higher for the higher silicon level. This resulted in better fluidity with an increase in silicon level.

90 5.2. FLOW STOPPAGE MECHANISM 1) The maximum melt temperature curve along the flow length showed a plateau in each case. The microstructure revealed that the flow stoppage occurred at that area due to heavily interlocked PSC. The high latent heat of fusion of silicon resulted in the rise in the temperature, resulting in the plateau and subsequent re-flow of metal.

2) The distance of the interlocked PSC zone from the gate increased with an increase in silicon level due to an increase in the solidification range.

5.3. EFFECT OF MODIFIERS ON EUTECTIC SILICON AND PSC MORPHOLOGY 1) Microstructure of the flow tips of unmodified and modified alloys primarily consisted of eutectic silicon in the aluminum matrix with a few PSC. This PSC depleted zone was the result of metal flow after flow stoppage due to interlocked PSC. Only a few PSC that drifted along with the flow could be seen at the tip. The gate area, however, consisted of both PSC and eutectic silicon in the aluminum matrix.

2) The PSC size increased with an increase in silicon level and in small amounts with increases in superheat. The mean PSC size in the 6.25 mm section was 4% – 12% smaller than that in the 12.50 mm section. This is attributed to an increased cooling rate in the thin section (39.4ºC/min) relative to that in thick section (26.2ºC/min).

3) PSC floatation was very significant on the 31.75 mm high runner in the Al-22%Si alloy. However, it was not observed in the 6.25 mm thick spiral and was not clearly distinguishable in the 12.50 mm thick spiral. PSC floatation was also not observed for Al-14%Si and Al-18%Si.

4) The severity of PSC floatation increased with an increase in superheat of the alloy and the casting thickness. It was also observed that the bigger size PSC floated to the top faster than the smaller ones due to higher buoyancy on the bigger PSC. This also resulted in an increase in hardness on the upper surface of the casting at the gate area.

91 5) Larger PSC were observed to be attached to the wall and the growth was in the opposite direction of heat flow. This observation shows that the PSC grew in dendrite-shape.

6) CMA modified the eutectic silicon from a needle-like shape to a finer shape but did not affect the PSC. With an increase in CMA additions, the level of modification increased from needle-like to lamellar to fine fibrous eutectic silicon. However, increases in the CMA addition led to an increased amount of Mg rich phases precipitations. Also, CMA did not affect the PSC and hence did not affect the PSC floatation.

7) Phosphorus did not affect the eutectic silicon but modified the PSC by reducing the size and changing the shape. This led to reduced floatation of PSC due to less buoyancy on the small size PSC.

8) Surface hardness of the alloys at the gate and flow tip increased with an increase in silicon level as a result of an increased number of PSC. Addition of CMA increased the surface hardness at the flow tip (eutectic silicon rich area), while phosphorus had an impact on the surface hardness at the gate (PSC rich area). Due to the presence of very hard PSC, the hardness at the gate was about 50% more.

92 SUGGESTION FOR FUTURE WORK

1. Study the floatation and modification of PSC by increasing the phosphorus content in the alloy. 2. Study the combined effect of CMA and phosphorus using a small size melt, where the chances of agglomeration of additives will be minimized. 3. Study the effect of CMA on the grain size of the alloys with various levels of CMA. 4. Study the effect of solution heat treatment on CMA modified alloys with 2.1% CMA.

Dissolution of Mg2Si may lead to further strengthening of the alloys. 5. Study the effect of CMA on eutectic reaction and fraction solid at various reactions. 6. Study the effect of CMA and phosphorus on the dendrite coherency of hypereutectic Al- Si alloys. 7. Study the effect of the modifiers on the tensile strength of the alloys. 8. Carry out detailed thermal analysis to predict the changes in the reaction temperatures as a result of modifier additions.

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96 38. Li, Z., Samuel, F., Samuel, A., Ravindran, C., Valtierra, S., Doty, H., “Factors Affecting Dissolution of CuAl2 Phase in 319 Alloys”, AFS Transactions, Vol.111, pp.1-14 (2003) Best Paper Award 39. Dahle, A.K., Nogita, K., McDonald, S.D., Dinnis, C., Luc, L., “Eutectic modification and microstructure development in Al–Si Alloys”, Materials Science and Engineering A 413–414, pp. 243–248 (2005) 40. Xu, C.L., Jiang, Q.C., “Morphologies of primary silicon in hypereutectic Al–Si alloys with melt overheating temperature and cooling rate”, Materials Science and Engineering A 437, pp. 451–455 (2006) 41. Loper. Jr, C.R., Cho, J.I., Hur, B.Y., Groteke, D.E., “Floatation of Primary Silicon in Under-Refined Melts”, AFS Transactions, pp. 243-249, (1998) 42. Donahue, R.J., Hesterberg, W.G., Cleary, T.M., “Hypereutectic Aluminum Silicon Alloy”, US Patent number 4969428, (Apr. 14 1989) 43. Djurdjevic, M., Gallo, P., Jiang, H., Sokolowski, J. H., “Evaluations of Strontium Fading in the 319 Aluminum Alloy using Thermal Analysis”, AFS Transactions, vol. 108, pp. (2000) 44. Djurdjevic, M. B., Kierkus, W. T., Byczynski, G. E., Stockwell, T. J., Sokolowski, J. H., "Modeling of Fraction Solid for the 319 Aluminum Alloy", AFS Transactions, vol. 107, pp. 173-179 (1999) 45. Hong, Soon-Jik, Suryanarayana, C., "Mechanical Properties and Fracture Behavior of an Ultrafine-Grained Al-20 Wt Pct Si Alloy", Metallurgical and Materials Transactions, vol. 36A, pp. 715-723 (Mar 2005) 46. Hengcheng Liao, Yu Sun, Guoxiong Sun, “Correlation Between Mechanical Properties and Amount of Dendritic -Al Phase in as-Cast Near-Eutectic Al–11.6% Si Alloys Modified with Strontium”, Materials Science and Engineering A335, pp. 62–66 (2002) 47. Pennors, A., Samuel, A.M., Samuel, F.H., Doty, H.W., “Precipitation of β – Al5FeSi Iron Intermetallic in Al-6%Si-3.5%Cu (319) Type Alloys: Role of Sr and P,” AFS Transactions, pp. 251-264 (1998) 48. Vasisht, S., Tsvetkov, A., Troczynski, T., Ceramic Master Alloy for Grain Refinement and Modification of Aluminum Alloys, AUTO21 conference, Barrie, Ontario (2006)

97 49. Samuel, F.H., Ouellet, P., Samuel, A.M., Doty, H.W., “Effect of Mg and Sr Additions on the Formation of Intermetallics in Al-6 Wt Pct Si-3.5 Wt Pct Cu- (0.45) to (0.8) Wt Pct Fe 319-Type Alloys,” Metallurgical and Materials Transactions A, vol. 29A, pp. 2871-2884 (Dec 1998) 50. Gowri, S., Samuel, F., ""Effect of Mg on the Solidification Behavior of Two Al-Si- Cu-Fe-Mg (380) Die Casting Alloys," AFS Transactions, vol. 101, pp.611-617 (1993) 51. Murali, S., Raman, K., Murthy, K.S.S., ""Effect of Magnesium, Iron (impurity) and Cooling Rates on the Fracture Toughness of Al-7 Si-0.3 Mg Casting Alloy," Material Science and Engineering, A1 51, pp. 1-10 (1992) 52. Cáceres, C. H., Sokolowski, J. H., Gallo, P., "Effect of Ageing and Mg Content on the Quality Index of Two Model Al-Cu-Si-Mg Alloys", Materials Science and Engineering A, pp. 53-61 (ISSN 0921-5093) (1999) 53. Ho, C.R., Cantor, B., “Heterogeneous Nucleation of Solidification of Si in Al-Si and Al-Si-P Alloys,” Acta Metallurgica et Materialia, vol. 43, No. 8, pp. 3231-3246 (1995) 54. Kyffin, W. J., Rainforth, W. M., Jones, H., "Effect of phosphorus additions on the spacing between primary silicon particles in a Bridgman solidified hypereutectic Al-Si alloy", Journal of Materials Science 36, pp. 2667 – 2672 (2001) 55. Sterner Rainer, R., “Aluminum Silicon Alloy with a Phosphorus Content of 0.001 to 0.1%”, US patent 1940922, (26 Dec. 1933) 56. Onitsch–Modl, E. M., in proceedings 1st world metallurgical congress, edited W. M. Baldwin, Jr. (American Society for Metals, Cleveland, 1952) pp. 325–337 (1951) 57. Mascr´e, C., Foundry Trade Journal, No. 94, pg. 725 (1953) 58. Loper. Jr., C.R., Cho, J.I., “Influence of trace amounts of Phosphorus in Al casting alloys-A review of literature; AFS Transactions, vol. 108, pp. 667-672 (2000) 59. Suzuki, T., Montani, M., Oshiro, N., “Effect of Phosphorus on Mechanical Properties of Al-Si Alloy”, Keikinzoku Gakkan Taikan, Koen Gaiyo, vol. 88, pp. 13-14 (1995) 60. Kato, E., Nomutra, H., Oshiro, N., “Microstructure of High Purity Al-Si Eutectic Alloys Containing Minute Amounts of Phosphorus”, Journal of Japan Institute of Light Metals, vol. 46, No. 8, pp. 377-382 (1996)

98 61. Oshiro, N., Suzuki, T., Kato, E., “The Modification of Microstructures and Mechanical Properties of JIS AC3A Alloys by Dephosphorizing”, Journal of Japan Institute of Light Metals, vol. 47, No. 6, pp. 341-346 (1996) 62. Xiangfa Liu, Yuying Wu, Xiufang Bian, “The nucleation sites of primary Si in Al–Si alloys after addition of boron and phosphorus”, Journal of Alloys and Compounds 391, pp. 90–94 (2005) 63. Bates, A.P., Calvert, D.S., “Refinement and foundry characteristics of hypereutectic Aluminum- Silicon alloys,” The British Foundryman, vol. 59, pp. 119-33, (1966) 64. Sulzer, J., “How to grain refine high silicon aluminum alloys,” modern casting, vol. 39, pp. 38-43 (1961) 65. Bichler, L., Ravindran, C., “A Study of EPS Bead Morphology and its Dependence on Bead Fusion”, AFS Transactions, paper 07-043 (2007) 66. Jue, B., Ravindran, C., Karpynczyk, J., “Permeability of Refractory Coating in EPC Process”, AFS Transactions, vol. 101, pp. 955-959, (1993) 67. Venkatramani, R., Ravindran, C., “Effects of Coating Thickness and Pouring temperature on Thermal Response in Lost Foam Casting”, AFS Transactions, vol. 104, pp. 281-290 (1996) 68. Ravindran, C., Jue, B., Karpynczyk, J., “Study of Effect of Vibration on Permeability of Unbonded Sand in EPC”, AFS Transactions, vol. 102, pp. 915-920 (1994) 69. Li, Kun-Dar, E. Chang, "Explanation of the Porosity Distribution in A206 Aluminum Alloy Casting", AFS Transactions, vol. 111, paper# 03-108 (2003) 70. Savitzky, M. Golay, “Smoothening and Differentiation of Data by Simplified Least Squares Procedure”, Analytical Chemistry, vol. 36, pp. 1627-1639 (1964) 71. Kierkus, W. T., Sokolowski, J. H., "Recent Advances in Cooling Curve Analysis: A New Method of Determining the 'Base Line' Equation", AFS Transactions, vol. 107, 161-167 (1999) 72. Djurdjevic, M., Stockwell, T., Sokolowski, J., “Effect of Sr on the Microstructure of Al-Si and Al-Cu Rich Eutectics in the 319 Alloy”, International Journal of Cast Metals Research, vol. 12, pp. 67-73 (1999) 73. Simpson, R., Ravindran, C., “Control of Microstructure in Lost Foam Casting Using Chills”, AFS Transactions, vol. 105, pp. 787-793 (1997)

99 74. Poniewierski, Z., Krystalizacja, struktura I wtasności siluminów (Crystallization, structure and properties of silumins), Wyd. Naukowo-Techniczne, Warsaw, (1989) (in Polish) as seen in “Aluminum-Silicon Casting Alloys: Atlas of Microfractographs” (#06993G), Chapter 1, “Introduction to Aluminum-Silicon Casting Alloys”, ASM International (2004) 75. Mondolfo, L.F., Aluminum alloys: Structure and properties, Butterworth, London- Boston, (1976) 76. Hatch, J.E., “Aluminum – Properties and Physical Metallurgy, American Society for Metals, pp. 30-31 (1984) 77. Russo, E. di, “The Atlas of Microstructures of Aluminum Casting Alloys”, Edimet, Brescia (1993) 78. Li, S.P., Zhao, S.X., Pan, M.X., Zhao, D.Q., Chen, X.C., Barabash, O. M., "Eutectic Reaction and Microstructural Characteristics of Al(Li)-Mg2Si Alloys", Journal of Materials Science, vol. 36, pp. 1569 – 1575 (2001) 79. Zhang, J., Fan, Z., Wang, Y. Q., Zhou, B. L. "Effect of Cooling Rate on the

Microstructure of Hypereutectic Al-Mg2Si Alloys", Journal of Materials Science Letters, vol. 19, pp. 1825 – 1828 (2000)

100 APPENDIX

Table A.3.1. Composition of Al-14%Si alloys with additives

Elements Al-14%Si Al-14%Si + CMA Al-14%Si + Phosphorus

Al 81.2 ± 0.4 80.0 ± 0.3 81.0 ± 0.4

Si 14.0 ± 0.4 14.0 ± 0.4 14.0 ± 0.4

Cu 2.95 ± 0.07 2.95 ± 0.09 2.95 ± 0.09

P 0.0014 0.0029 0.0138

Mg 0.0400 0.7340 0.0360

Sr <0.0002 <0.0002 <0.0002

Fe 0.5880 0.8327 0.7415

Mn 0.2342 0.4099 0.3423

Cr 0.0432 0.0737 0.0649

Pb 0.0625 0.0526 0.0438

Ni 0.0550 0.0419 0.0435

Zn 0.7038 0.7496 0.6869

Ti 0.1125 0.1060 0.1104

101

Table A.3.2. Composition of Al-18%Si alloys with additives

Elements Al-18%Si Al-18%Si + CMA Al-18%Si + Phosphorus

Al 77.7 ± 0.4 77.0 ± 0.4 77.5 ± 0.4

Si 18.0 ± 0.2 18.0 ± 0.2 18.0 ± 0.2

Cu 2.75 ± 0.5 2.75 ± 0.3 2.75 ± 0.3

P 0.0019 0.0017 0.0109

Mg 0.0350 0.7340 0.0630

Sr <0.0002 <0.0002 <0.0002

Fe 0.4754 0.8752 0.7806

Mn 0.2066 0.3395 0.2889

Cr 0.0383 0.0566 0.0486

Pb 0.0129 0.0100 0.0165

Ni 0.0801 0.0590 0.0618

Zn 0.6698 0.5803 0.6109

Ti 0.1161 0.0873 0.0928

102

Table A.3.3. Composition of Al-22%Si alloys with additives

Elements Al-18%Si Al-18%Si + CMA Al-18%Si + Phosphorus

Al 74.2 ± 0.5 73.5 ± 0.4 74.0 ± 0.6

Si 22.0 ± 0.2 22.0 ± 0.2 22.0 ± 0.2

Cu 2.45 ± 0.7 2.45 ± 0.4 2.45 ± 0.4

P 0.0022 0.0023 0.0105

Mg 0.0490 0.6290 0.0520

Sr <0.0002 <0.0002 <0.0002

Fe 0.5178 0.6994 0.6476

Mn 0.1738 0.2097 0.2055

Cr 0.0304 0.0360 0.0449

Pb 0.0138 0.0166 0.0169

Ni 0.0269 0.0366 0.0368

Zn 0.4286 0.5258 0.4780

Ti 0.0894 0.0795 0.0807

103

Figure A.4.1. Heating and cooling curves with the first derivative curve for Al-14%Si

104

Figure A.4.2. Heating and cooling curves with the first derivative curve for Al-18%Si

105

Figure A.4.3. Heating and cooling curves with the first derivative curve for Al-22%Si

106

Table A5. List of materials and suppliers

Sl. No. Material Purpose From Telephone

Intended to prepare Tim Moses, House of 416-421-1572 X27 (Work) 1 Pure Silicon Al -50%Si master alloy metals 519-658-3056 (cell)

Lucas Milhaupt (A Addition of Phosphorus 414-769-6000 2 Cu-P refiners Handy Harman as grain refiner www.lucasmilhaupt.com Company)

Al-50%Si master To vary the %Si in the David Ogden, 3 1-800-523-8457 X 402 (work) alloy alloy for experiments K. B. Alloys

Use the composite Tom Troczynski, 604-822-2612 CMA (refiner and 4 master alloy as grain University of British email: modifier) refiner & modifier Columbia [email protected]

Base alloy for Sailesh, 5 A319 Al-Si alloy 519-827-9423 experiments Linamar, Guelph

107 Curriculum Vitae

Name: Achinta Haldar Education: M.A.Sc. Mechanical Engineering (2007) Thesis: A Study on Fluidity, Flow Stoppage and Silicon Modification of Hypereutectic Aluminum-Silicon Alloys in Lost Foam Casting B.Sc. Metallurgical Engineering (1990) Thesis: Effect of Heat Treatment Processes of Fatigue Life of Low Carbon Steel

Work Experience: Plant Engineer, Concord Heat Treat Ltd. Ontario, Canada, (2004-2005) Foundry Technician, K. P. Bronze Ltd, Ontario, Canada, (2003-2004) Quality Inspector, ABC Plastic Molding, Ontario, Canada, (2002-2003) Process Engineer, Steel Authority of India Ltd. Rourkela, India (1991-2002) Engineer Trainee, Hindusthan Development Corporation, (1990-1991)

Professional Memberships: Professional Engineers of Ontario (PEO) American Society of Metals (ASM) (International) American Foundry Society (AFS) Association for Iron and Steel Technology (AIST) Foundry Education Foundation (FEF) Canadian Institute of Mining, Metallurgy and Petroleum (CIM) American Society for Quality (ASQ)

Awards & Scholarships: Ryerson Graduate Scholarships for 2005-06 and 2006-07 Merit scholarship from the American Foundry Society, Ontario chapter, 2006 “Past President’s Award” from Canadian Foundry Association, 2006

Publications: 1) A. Haldar and C. Ravindran, “Effect of Superheat and Section Thickness on Flow Characteristics for Thin Section LFC of Hypereutectic Al-Si Alloys”, Metsoc COM 2007, International Symposium of Light Metals, Toronto. 2) A. Haldar and C. Ravindran, “Study of Flow Stoppage in Hypereutectic Al-Si Alloys in the Lost Foam Casting Process”, Metsoc COM 2007, International Symposium of Light Metals, Toronto. 3) A. Haldar and C. Ravindran, “Effect of CMA and Phosphorus on Silicon Modification and Floatation in LFC”, AUTO21 Conference, Windsor, 2007 4) A. Haldar and C. Ravindran, “A Study on Fluidity and Silicon Modification in Lost Foam Casting of Hypereutectic Al-Si Alloys”, AFS Transaction, 2008 (in preparation). 5) A. Haldar and C. Ravindran, “Effect of Superheat and Silicon Level on Flow Characteristics and Silicon Morphology on Thin Section LFC of Hypereutectic Al-Si Alloys”, Materials Characterization, 2008, (in preparation).

108