quantification of iron in alsi foundry...

QUANTIFICATION OF IRON IN ALSI FOUNDRY ALLOYS VIA THERMAL ANALYSIS

Robert Ian Mackay

Department of Mining & Metallurgicai Engineering McGa University

Montreal, Quebec, Canada

A Thesis subrnitted to the Faculty of Graduate Studies & Research

in partial fulfilment of the requirements of the degree of

Master of Engineering

R. Mackay, 1996

National tibrary 1*1 of Canada Bibliothèque nationale du Canada

Acquisitions and Acquisitions et Bibliographie Services services bibliographiques

395 Wellington Street 395, nie Wellington Ottawa ON K I A ON4 Ottawa ON K I A ON4 Canada Canada

The author has granted a non- L'auteur a accordé une licence non exclusive licence dowing the exclusive permettant à la National Library of Canada to Bibliothèque nationale du Canada de reproduce, loan, distribute or sell reproduire, prêter, distribuer ou copies of this thesis in microfom, vendre des copies de cette thèse sous paper or electronic formats. ia forme de microfiche/nlm, de

reproduction sur papier ou sur format électronique.

The author retains ownership of the L'auteur conserve la propriété du copyright in this thesis. Neither the droit d'auteur qui protège cette thèse. thesis nor substantid extracts fiom it Ni la thèse ni des extraits substantiels may be printed or otheMrise de celle-ci ne doivent être imprimés reproduced without the author's ou autrement reproduits sans son permission. autorisation.

To my Parents Dr. Alward & Yvonne Mackay

Abstract

Iron content in aiuminum-silicon casting ailoys, which contribute to the formation of

AisFeSi intennetallic, cm be very demmental to the physical properties of the final cast

part. Because of the brittle and hard nature of the AlPeSi intermetallic, machinhg of

casting parts can be difficult and wstly to the foundry due to the ne& to increase

operathg horse power of the machining apparatus and reduced tool life.

Thermal analysis could provide a cost effective and reliable rnethod to quantify

the iron content of the alloy melt before the casting process is perfonned. The formation

of the A1,FeSi intermetallic can be resolvable on the cooling curves for alurninum-sificon

alloys if the iron content is equal to or greater than 0.6%wt when the cooling rate is

O. 1O"Clsec. As the iron content gradually increases, the formation temperature of Al,FeSi

increases and this results in an increase in the duration of the GFeSi thermal anomaly.

A time based parameter associated with the A15FeSi thermal anomaly is also used to

quanti9 the Fe content. Time parameten can be very accurate if the melt volume and

heat extraction for the solidifying thermal analysis sample are strictly controlied.

Results of Fe quantification via apparent time parameter of the A&FeSi

c r y s ~ t i o n for 356,319 and 413 dioys using thermal analysis has been wmpleted for

this thesis. The results show that a proportionality between percent area of A1,FeSi found

in the microstructure using the Leco image analyzer, and to the percent Fe determined

b y spectrochemical analysis , exists with the apparmt ti me parameter associated with the

A1,FeSi thermal anomaiy. It is also shown that the duration of the eutectic growth, as

determined fom cooling curves, is effeaed by the presence of iron. This indicates that

time parameters for the AlPeSi could be used by foundrymen to quanti@ Fe content of

their aluminum-silicon melts.

En contribuant la formation d'interm~taiiiques AisFeSi, le fer présent dans les

aiiiages de fonderie d'aluminium-silicium peut &me préjudiciable aux propriktés physiques

du produit coule final. La nature dure et fragile des intenn~talliques AISFeSi peut rendre

l'usinage des produits coulés, difficile et coûteux pour la fonderie A cause de la nécessité

d'augmenter la puissance de l'appareii d'usinage et de la durée de vie réduite de l'outil.

L'analyse thermique pourrait fournir une mtthode fiable et peu chère pour

quantifier la teneur en fer de l'alliage avant que celui-ci soit coulé. La formation des

interrnetalliques AISFeSi peut être observée sur les courbes de refroidissement des alliages

d'aluminium-silicium si la teneur en fer est egale ou supérieure 0.6 pct en poids et pour

une vitesse de refroidissement de O.lO°C/sec. Lorsque la teneur en fer augmente

graduellement, la température de formation de A1,FeSi augmente et ceci entraine une

augmentation de la durée de l'anomalie thermale associée à Al&Si. Un parametre basé

sur le temps, associé B l'anomalie thermale de AlsFeSi a aussi W utilisé pour quantifier

la teneur en fer. Les paramktres de temps peuvent être tri3 précis si le volume de metal

fondu et l'extraction de chaleur pour l'échantillon d'analyse thermique se solidifiant sont

strictement controllés.

La quantification du fer via un pararn&tre de temps apparent de la cristallisation

de A1,FeSi a été faite pour les alliages 356, 319 et 413 en utilisant l'analyse thermique.

Les résultats montrent qu'une proportionalité existe entre le pourcentage d'aire de

AIsFeSi trouvé dans la microstructure et le pourcentage de Fe déterminé par analyse

spectrom&rique. Ii a aussi CtC montre que la durée de la croissance eutectique, telle que

determin& A partir des courbes de refroidissement, est affectée par la présence de fer.

Ceci indique que les paramktres de temps pour la phase A1,FeSi pourraient &re utilisés

par les fondeurs pour quantifier la teneur en fer des coulées d'aluminium-silicium.

Acknowledgements

First I would Like to thank my supervisor Dr. John Gruzieski for taking me as a graduate

student and supporting me with a stipend fkom his grant fkom the Natural Sciences and

Engineering Research Council of Canada (NSERC). 1 would also like to thank Dr.

Gnitleski for the suggestion of the thesis topic and supervising the work completed for

this Master of Engineering degree.

1 am grateful to my parents Dr. Alward and Mrs. Yvonne Mackay for the

financial and emotionai support that was so critical during my move corn St. John's,

Newfoundland to Montreal, and during my stay in Montreal to pursue this M.Eng

degree.

1 would like to thank Dr. Florence Paray for the supervision and advice given to

me in the foundry laboratory, emission spectrometer laboratory, Lem image analyzer and

metallographic laboratory .

I would like to thank Dr. Joe Hodych & Ray Patzold of the Department of Earth

Science, Mernorial University of New foundland , S t. John's, Newfoundland . Dr. Jerry H.

Sokolowski, Alan Esseltine, Heather Mckechnie & Tony Graci of the Department of

Mechanical Engineering, Windsor University, Windsor, Ontario. Charles Antony

Bhaskaran and David Sparkman of Foundry Information Systems, New Castle, Indiana.

Helen Campbell, Priti Wanjara, Leslie Ederer, Chris Carozza, Monique Riendeau,

Marlene Gray, Daryuosh Emadi, Robert Paquette, James Ansen, Tom Ledermen, David

Gloria, Chito Edovas, Saeed Shabestari, Ahmed Abdollahi, Musbah Mahfoud and

Ramani Sankaranarayanan of the Department of Mining and Metailurgical Engineering,

McGill University, Montreal, Quebec.

Table of Contents

. . Abstract . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . u ... . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . R6sum6 m

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Acknowledgements iv . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Table of Contents v

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . List of Figures sz

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . List of Tables .xii

Chapter One : Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1 .O Introduction 1

1.1 Iron in Alurninum Ailoys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 1.3 Future of Aiuminum Ailoy Casting . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 1.4 Objective of Thesis Project . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3

Chapter Two : Literature Review . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2 Sources of iron in Aluminum Alloys 4

2.2.0 High Pressure Die casting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4 2.2.1 Low Pressure Permeant Mold casting . . . . . . . . . . . . . . . . . . . . . . 6 2.2.2 Aluminum Recycling . . . . . . . . . . . . . . . . . . . . ... . . . . . . . . . . 6

2.1 Fe-bearing intennetallics in Alurninum-Silicon Alloys . . . . . . . . . . . . . . . . 7 2.1.1 Effects of B(AlFeSi) Phase on

PhysicaI Properties of AI-Si AUoys . . . . . . . . . . . . . . . . . . . . . . . . 10 2.1.1 .O General Physicd Properties . . . . . . . . . . . . . . . . . . . . . . . 10 2.1.1.1 Effect of Fe-intermetallics

on Corrosion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .10 2.1.1.2 Effect of Fe-intermetallics

on Eutectic solidification and Modification . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11

2.1.1.3 Effst of Fe content on Grain Size . . . . . . . . . . . . . . . . . . . 11 2.1.1.4 Effect of Fe content on

. . . . . . . . . . . . . . . . . . . . . . . . . . Dendrite A m Spacing - 11 2.1.1.5 Effect of Fe-intermetallics on

Post-casting Machinhg . . . . . . . . . . . . . . . . . . . . . . . . . .ll 2.1.1.6 Effect of Fe-in terrnetaliics on

. . . . . . . . . . . . . . . . . . . . Elevated Temperature Properties 12 2.2.2 Methodology For Correcting the Deleterious

Effects of P(A1Fe-i) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .12 2.2.2.0 Master Alloy Addition . . . . . . . . . . . . . . . . . . . . . . . . . . . 12 2.2.2.1 Liquid Alloy Superheat . . . . . . . . . . . . . . . . . . . . . . . . . . 14 2.2.2.2 Cooling Rate . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . -14 2.2.2.3 Solution Heat Treatment . . . . . . . . . . . . . . . . . . . . . . . . . 15 2.2.2.4 Dilution . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . -15

vi

Page

Chapter Thme : Thermal Analysis of Aluminum-Siüeon Aïîoys 3.0 Thermal Analysis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16

3.0.1 Introduction to Basic Concepts . . . . . . . . . . . . . . . . . . . . . . . . . . 16 3.0.2 Thermal Analysis of Typicai M-Si Moys . . . . . . . . . . . . . . . . . . . 17

3.1 The Therrnal Anaiysis of Minor Reactions . . . . . . . . . . . . . . . . . . . . . . 18 3.1.0 B(AiFeSi) Reactions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 18 3.1.2 Other Post-eutectic Reactions . . . . . . . . . . . . . . . . . . . . . . . . . . . 19 3.1.3 Thermal Analysis Parameten . . . . . . . . . . . . . . . . . . . . . . . . . . . 19 3.1.4 Limitations of Thermal Analysis . . . . . . . . . . . . . . . . . . . . . . . . . 22

Chapter Four : Experimental Methodology 4.0 Chemistry & Microstructure of Ai-Si Ailoys Studied . . . . . . . . . . . . . . . . 23

4.0.1 3 19 (Al-Si-Cu) Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 23 4.0.2 356 (Al-Si-Mg) Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 24 4.0.3 413 ( A M ) Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25

4.1 Experimental Foundry Procedure . . . . . . . . . . . . . . . . . . . . . . . . . . . . 29 4.2 Experimentai Setup for Thermal Analysis . . . . . . . . . . . . . . . . . . . . . . . 29 4.3 Post Foundry Experirnentai Analysis . . . . . . . . . . . . . . . . . . . . . . . . . . 31

4.3.1 Spectrochemical Analysis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 31 4.3.2.1 Polishing of Aluminum Alioy Samples . . . . . . . . . . . . . . . . . 32 4.3.2.2 Image Analysis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 32

Chapter Five : Results of Thermal Analysis 5.0 Spectrochemical Analysis

of Aluminum-Silicon Aiioys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . -33 5.1 Results of Thennd Analysis & Associated Microstructure . . . . . . . . . . . . . 34

5.1.1 A356 Alloy (O.lO°C/sec & 0.45"Clsec) . . . . . . . . . . . . . . . . . . . . 34 5.1.2 319 Alloy (O.lO°C/sec) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 42 5.1.3 A413 Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 42

5.2 Quantification of Fe-bearing Intermetallics . . . . . . . . . . . . . . . . . . . . . . 53 5.2.1 AisFeSi Thermal Signature . . . . . . . . . . . . . . . . . . . . . . . . . . . . 53 5.2.2 QuantiQing Fe via Thermal Andysis . . . . . . . . . . . . . . . . . . . . . . 53

5 .2.2.1 Temperature Method . . . . . . . . . . . . . . . . . . . . . . . . . . . . 54 5.2.2.2 Time Method . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 55 5.2.2.3 Time & Temperature Method for

Eutectic Solidification . . . . . . . . . . . . . . . . . . . . . . . . . . . . 55 5.3 Tabular Results of Thermal Analysis Parameters . . . . . . . . . . . . . . . . . . . 63

5.3.1 General Trends Observed in Tabled Results . . . . . . . . . . . . . . . . . . 70 5.4 Image Analysis Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 70

5.4.1 Image Anaiysis of A356 AUoy . . . . . . . . . . . . . . . . . . . . . . . . . . 71 5.4.2 Image Analysis of 319.2 Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . 71

. . . . . . . . . . . . . . . . . . . . . . . . . . 5.4.3 Image Analysis of A413 AUoy 71 5.5 Cornparison of Formation Temperature & Apparent Time

Parameter for Quantifying Iron content of . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Aluminurn-Silicon AUoys -75

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.5.1A356Alloy 75 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.5.2 319.2 Alloy 81 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.5.3A413A.h~. 81

. . . . . . . . . . . . . . . . . . . . . . . 5.6 Effect of Iron on Eutectic Tirne Duration 82 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.6.1 A356 Aiioy 82 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.6.2 A413 Alloy 82 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.6.3 319.2 AUoy 83

. . . . . . . . . . . . . . . . . . . . . . . . . 5.6.4 Eutectic Formation Temperature 83

Chapter Six : Discussion of Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.0 Discussion of Results 89

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.0.1 Hypoeutectic Alloys 89 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.0.2 Eutectic Alloys -93

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.1 Quality of Denvative Curves -95

Chapter Seven : Conclusions & Future Work . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.0 Conclusions -97 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.1 Future Work -99

List of Figures

. . . . . . . . . . . . . . . Figure 2.Oa Schernatic of High Pressure Die Casting Unit 5 . . . . . . . . . . . . . . . . . . Figure 2.0b Schematic of low-Pressure Casting Unit 5

. . . . . . . . . . . . . . . . Figure 2 . la Liquidus Surface of Al-Fe-Si Phase Diagram 8 . . . . . . . . . . . . . . . . . . . . . . Figure 2 lb Al-Fe Pseudo-Binary Phase Diagram 8

. . . . . . . . . Figure 2.2 Twin Plane Re-entrant Edge (TPRE) Growth Mechanism 9 Figure 2.3 Simplified phase diagram of the Al-Fe-Si-Mn

. . . . . . . . . . . . . . . . . . . . . . . . . system with constant Mn levels 13

Figure 3.0 Alurninum-Silicon Cooling Curve & Derivative . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Cuwe (319.2 alloy) 17

Figure 3.1 Aluminum-Silicon Cwling Curve & Denvative Curve (3 19.2 alloy) with high iron (0.8 % wt) . . . . . . . . . . . . . . . . 18

. . . . . . . . . . . . . . . . . . . . . . . . . . Figure 3.2 Thermal Analysis Parameters 20 Figure 3.3 Thermal Analysis Time Parameters . . . . . . . . . . . . . . . . . . . . . . 20 Figure 3.4 Thermal Analysis of a near eutectic ductile cast iron . . . . . . . . . . . . 22

Figure 4.0a Microstructure of a 319 alloy . . . . . . . . . . . . . . . . . . . . . . . . . 27 Figure 4.0b Microstructure of a 356 alloy . . . . . . . . . . . . . . . . . . . . . . . . . 27 Figure 4 . 0 ~ Microstructure of a 413 alloy . . . . . . . . . . . . . . . . . . . . . . . . . 28 Figure 4.1 Thermal Analysis Setup . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 30

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . Figure 4.2a Stainless Steel Crucible 30 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Figure 4.2b Sand Mold Setup 30

Figure 5 . l a 356 Cooling Curve & Associated Denvative (Fe content =0.54% wt. cooling rate = O.lO°C/sec) . . . . . . . . . . . 36

Figure 5 . lb Microstmcture of associated 356 alloy in fig.5. la . . . . . . . . . . . . . 36 Figure 5.2a 356 Cooling Curve & Associated Denvative

(Fe content =0.98% wt. cooling rate = O.lO°C/sec) . . . . . . . . . . 37 Figure 5.2b Microstructure of associated 356 aUoy in fig.5.2a . . . . . . . . . . . . . 37 Figure 5.3a 356 Cooling Cuve & Associated Derivative

(Fe content =1.25% wt. cooling rate = O.lO°C/sec) . . . . . . . . . . 38 Figure 5.3b Microstructure of associated 356 ailoy in fig.5.3a . . . . . . . . . . . . . 38 Figure 5.4a 356 Cooling Curve & Associated Derbative

(Fe content =0.75 % wt, cooling rate = 0.45'C/sec) . . . . . . . . . . 39 Figure S.4b Microstructure of associatcd 356 aiioy in fig.5.4a . . . . . . . . . . . . 39 Figure 5.5a 356 Cooling Curve & Associated Derivative

(Fe content =0.90 % wt, cooling rate = 0.45 "Clsec) . . . . . . . . . . 40 Figure 5.5b Microstructure of associated 356 aUoy in fig.5.5a . . . . . . . . . . . . 40 Figure 5.6a 356 Cooling Curve & Associated Derivative

(Fe content = 1.40% wt, cooling rate = 0.4S°C/sec) . . . . . . . . . . 41 Figure S.6b Microstnicture of associated 356 aUoy in fig.5.6a . . . . . . . . . . . . 41

Page Figure 5.7a 3 19.2 Cooling Cuve & Associated Derivative (trial one)

. (Fe content =0.36% wt. cooling rate = O 10°C/sec) . . . . . . . . . . . 43 . . . . . . . . . . . . Figure 5.7b Microstnichire of associated 3 19.2 alloy in fig.5.7a 43

Figure 5.8a 3 19.2 Coohg C w e & Associated Derivative (trial one) (Fe content =0.90% wt, oooling rate = O . 10°Clsec) . . . . . . . . . . . 44

. . . . . . . . . . . . Figure S.8b Microstructure of associated 319.2 ailoy in fig.5. & 44 Figure 5.9a 319.2 C w h g Curve & Associated Derivative (trial one)

(Fe content =1.2% wt. cooling rate = O.lO°C/sec) . . . . . . . . . . . 45 . . . . . . . . . . . . Figure S.9b Microstnicture of associatecl 319.2 alloy in fig.5.9a 45

Figure 5.10a 3 19.2 Cooling Curve & Associated Denvative (trial hvo) . (Fe content =0.36% wt. cooling rate = O 10°C/sec) . . . . . . . . . . . 46

Figure 5 . lob Microstructure of associated 319.2 alloy in fig.5.10a . . . . . . . . . . 46 Figure 5.1 la 3 19.2 Cooling Curve & Associated Derivative (triai two)

(Fe content =0.92 % W. cooling rate = O . 10°C/sec) . . . . . . . . . . . 47 Figure 5.1 lb Microstructure of associated 319.2 ailoy in fig.5.l la . . . . . . . . . . 47 Figure 5.12a 319.2 Cooling Curve & Associatcd Derivative (hial two)

(Fe content = 1.25 % wt. cooling rate = O . 10°C/sec) . . . . . . . . . . . 48 Figure 5.12b Microstructure of associated 3 19.2 alloy in fig.5.12a . . . . . . . . . . 48 Figure 5.13a 4 13 Cooling Curve & Associated Derivative

. . . . . . . . . . . . . . . . . . . . . . . . . . . . F e content =0.92% wt) 49 Figure 5.13b Microstructure of associated 4 13 ailoy in fig .5 . 13a . . . . . . . . . . . 49 Figure 5.14a 4 13 Cooling Curve & Associated Derivative

(Fe content =1.4% wt) . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50 Figure 5.14b Microstructure of associated 413 alloy in fig.5.13a . . . . . . . . . . . 50 Figure 5.15a 413 Cooling Curve & Associated Derivative

(Fe content t2.295 wt) . . . . . . . . . . . . . . . . . . . . . . . . . . . . 51 Figure 5.15b Microstmcture of associated 413 aUoy in fig.5.13a . . . . . . . . . . . 51 Figure 5.16 Determination of apparent time parameters of the

AISFeSi intermetàllic for 356 & 319 . . . . . . . . . . . . . . . . . . . . . 57 Figure 5.17 Determination of apparent time parameter of the A15FeSi

intermetallic & pnmary aluminum of 413 alloy . . . . . . . . . . . . . . 58 Figure 5.18 Determination of eutectic solidification time for

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 356 & 319 dloys 3 9 Figure 5.19 Eutectic time determination for low iron 413 . . . . . . . . . . . . . . . . 60 Figure 5.20 Eutectic time determination for high iron 413 . . . . . . . . . . . . . . . 61 Figure S.2la Cooling curve with fourth derivative . . . . . . . . . . . . . . . . . . . . 62 Figure 5.21b Cornparison of second and fourth denvative . . . . . . . . . . . . . . . . 62 Figure 5.22 %Ara of @-phase in microstructure vs . apparent time

356 ailoy. reference cooling rate = O . 10°C/sec . . . . . . . . . . . . . . 72 Figure 5.23 %Ares of 8-phase in microstnicture vs . apparent time

356 ailoy. reference cooling rate = 0.45 "Clsec . . . . . . . . . . . . . . 72 Figure 5.24 %Area of &phase in microstructure vs . apparent time

3 19.2 (trial one). reference cooiing rate = O . 10°Clsec . . . . . . . . . 73

Page Figure 5.25 IArea of 8-phase in microstructure vs. apparent time

. . . . . . . . . 3 19.2 (triai two), reference cooling rate = O. 10°C/sec 73 Figure 5.26 %Axa of 8-phase in microstructure vs. apparent time

. . . . . . . . . . . . . . . . . . . . . . 413 aiioy, insulated cup method .74 Figure 5 -27 % Area of 8-phase in microstnicture vs. Eutectic Dmtion

. . . . . . . . . . . . . . . . . . . . . . 413 aiioy, insulated cup method .74 Figure 5.28 Formation temperature of &Phase vs. Iron Content

A356 alioy, rtference cooling rate = 0. 10°C/sa: . . . . . . . . . . . . . 76 Figure 5.29 Formation temperature of fi-Phase vs. Iron Content

A356 ailoy, reference cooling rate = O.4S0C/sec . . . . . . . . . . . . . 76 Figure 5.30 Apparent Time of fi-Phase Growth vs. Iron Content

. . . . . . . . . . . . . . A356 alloy reference cooling rate = O. 10°C/sec 77 Figure 5.3 1 Apparent Time of 6-Phase Growth vs. Iron Content

. . . . . . . . . . . . . A356 alloy, reference ml ing rate = 0.45OC/sec 77 Figure 5.32 Formation temperature of @-Phase vs. Iron Content

. . . . . . 319.2 ailoy (triai one), reference cooling rate = O. 10°C/sec 78 Figure 5.33 Formation temperature of &Phase vs. Iron Content

. . . . . . 3 19.2 alloy (trial two), reference cooling rate = 0.4S°C/sec 78 Figure 5.34 Apparent Time of 8-Phase Growth vs. Iron Content

. . . . . . 3 19.2 aUoy (trial one), reference oooling rate = O. 1 O O C/sec 79 Figure 5.35 Apparent Time of 8-Phase Growth vs. Iron Content

. . . . . . 319.2 ailoy (triai two), reference cooling rate = 0.45"C/sec 79 Figure 5.36 Formation temperature of 6-Phase vs. Iron Content

A413 ailoy, Insulated mold method . . . . . . . . . . . . . . . . . . . . . 80 Figure 5.37 Formation temperature of Primary Aluminum vs. Iron Content

A4 13 alloy . Insulated mold method . . . . . . . . . . . . . . . . . . . . . 80 Figure 5.38 Eutectic Duration vs. Iron Content

A356 ailoy, reference coolhg rate O. 10°C/sec . . . . . . . . . . . . . . 84 Figure 5.39 Eutectic Duration vs. Iron Content

A413 aUoy, insulated mold method . . . . . . . . . . . . . . . . . . . . . 84 Figure 5.40 Eutectic Duration vs. bon Content

. . . . . . . . 3 19.2 alloy (trail one), reference cooling rate 0. 10° C/sec 85 Figure 5.41 Eutectic Duration vs. Iron Content

3 19.2 ailoy (trial two) . reference cooling rate O. 1 O "Clsec . . . . . . . . 85 Figure 5.42 Eutectic Start & End Temperatures

3 19.2 ailoy (trial two), reference coolhg rate O. 10°C/sec . . . . . . . . 86 Figure 5.43 Eutectic S M & End Temperatures

A413 aiioy, insulated mold method . . . . . . . . . . . . . . . . . . . . .86 Figure 5.44 Detennination of the End of Eutectic Solidification

Local Maximum of the Second Derivative . . . . . . . . . . . . . . . . . 87 Figure 5.45 Detemination of the End of Eutectic Solidification

Maximum of the Fourth Derivative . . . . . . . . . . . . . . . . . 87

page Figure 6.0 Formation Temperature of j3-Phase vs. Iron Content -

. . . . . . . . . . . . . A356 dey, reference cooling rate = O. 10°C/sec 90 Figure 6.1 Formation Temperature of &Phase vs. Iron Content

. . . . . . . . . . . . . A356 aiioy, reference coohg rate = 0.45"C/sec 90 Figure 6.2 Formation Temperature of &Phase vs. Iron Content

319.2 aUoy (trail one), reference cooling rate = O.lO°C/sec . . . . . . 91 Figure 6.3 Formation Temperaîure of &Phase vs. Iron Content

319.2 aUoy (W two), reference cooling rate = O. 10°C/sec . . . . . . 91 . . . . . . . . . . . . . . . . Figure 6.4 Simplified Al-Fe-Si (Constant 0.25 % wt Mn) 94

List of Tables

. . . . . . . . . . . . Table 4.0 Aiuminum Associated Specifications for 3 19 Alloy 24

. . . . . . . . . . . . Table 4.1 Aluminum Associated Specifications for 356 Alloy 25

. . . . . . . . . . . . Table 4.2 Aluminum Associated Specifications for 413 Alloy 26

Table 5.1 Temperature and Time Based Parameters for A1,FeSi intermetallic in A356 Aiioy, Reference cooling rate = O.lO°C/sec . . . . . . . . . . . . 63

Table 5.2 Temperature and Time Based Parameters for AisFeSi intermetallic in A356 Alioy, Reference cooling rate = 0.4S°C/sec . . . . . . . . . . . . 64

Table 5.3 Temperature and Time Based Parameters for A15FeSi intermetallic in 3 19.2 Alloy (trail one), Reference cooling rate = 0.45 "Clsec . . . . 65

Table 5.4 Temperature and Time Based Parameters for A1,FeSi interrnetaliic in 3 19.2 Alloy (trial two), Reference cooling rate = 0.4S°C/sec . . . . . 66

Table 5.5 Eutectic Time & Temperature Parameters for 319 and 356 alloys . . . . 67 Table 5.6 Time parameters of the AisFeSi intermetallic, a-aiuminum &

Alurninum-Silicon Eutectic of 413 Woy . . . . . . . . . . . . . . . . . . . .68 Table 5.7 Temperatures parameters of the Al,FeSi intermetallic, a-aluminum &

Aluminum-Silicon Eutectic of 4 13 Alloy . . . . . . . . . . . . . . . . . . . . 69

CHAITER ONE : INTRODUCTION 1

Chapter One

Introduction

1.0 Introduction Aiuminum-siliwn cast alloys are finding numerous uses in the automotive and aerospace

industries and their use is expected to increase in these industries in the near future. The

reason for this increase is that these alloys have many desirable qualities such as a high

strength to weight ratio, good castability, excellent corrosion resistance, cosrnetic surface

quality, resistance to hot tearing, relatively good thermal conductivity, lower melting

temperatures, good machinability and good weldability. Another quality, besides light

weight, is that aluminum alloys are cornparatively easy to recycle, an important

consideration in this era of recyclability and environmental awarenessl!

Aluminum-silicon alloys are extremely important in the aluminum casting industry

since they make up 85% or more of the total aluminum cast parts produced2. These

aüoys usually contain between 5.5- 12 % silicon, and other wrnmon alloying elernents are

copper and magnesium3. Alterations of the microstructure, usually wmposed of a

dendritic aluminum matrix with an aluminum-silicon eutectic structure located in the

remaining spaces, can be done mainly via change of cooling rate, or by the addition of

certain master alloys to the melt prior to casting. Boron and titanium boride compounds

in the aluminum melt can increase the number of heterogenous nucleation sites and refme

the gain s i~e*~~ ' . The addition of sodium or strontium to the melt can modify the silicon

of the eutectic structure from an acicular to a fibrous rn~rphoiogf*~~*\ Addition of

strontium has also been shown to reduce heat treatment times9. Increase in cooling rate

can result in similar alterations to the rnicrostmcture. This variability of microstructure

for the same alloy type via cwling rate or master alloy additions allows aluminum-silicon

alloys to achieve a myriad of physical properties, giving them a wide range of uses in

the automotive, aerospace and other industry sectors.

ç- ONE : ~ O D U C ' r I O N 2

1.1 Iron in Aluminum AIloys In many aluminum-silicon casting alloys, iron nacts with durninum and silicon

to form a thmodynamicaliy stable phase having the stoichiometry of Al$%Sil*. As

discussed in the next chapter the morphology of this phase is in platelet form, but seen

as nde- l ike in a two-dimaisional metallographic c~oss-section~~. As a consequence,

one of the major difficulties with aluminum-silicon aüoys is that im, if present in large

enough quantities, can lead to this particular intermetallic having deleterious e f f m on

the physical properties of the cast part. In addition, these cornpounds are very hard with

the result that machining cast parts with a relatively high iron content can be difficult,

resulting in high casting finishing costs.

The size and number of such needle phases in the microstnicture of a casting is

difficult to determine in advance, and is one of the many reasons why customers using

alurninum cast parts rquire a casting or d e t y factor. This is costly for foundries and

increases the weight and size of the cast part.

1.3 Future of Aiuminurn Alloy Processing The quality of aluminum-silicon castings needs to improve if foundries are going

to be wmpetitive. Foundries need an accurate method of quantifjmg iron content on-line.

The only method used currently is spectrochemical analysis. In this process,

spectrochemical samples are poured and then taken off-line and analyzed. This operation

requires time and capital cost that smaller foundries canot fiord. A cost effective on-

line method to gauge the potential microstructure of the casting before pouring is clearly

needed. Such a method would be appealing for larger foundries to reduce operating costs

and maintain casting quality, while allowing smaüer foundries to improve casting quality

at a cost within their means. One possible approach to this problem is to use thermal

analysis. This particular quality control method is currently being used in many

aluminum-silicon foundries around the world to gauge the extent of grain refinement and

eutectic modification, however it has not been used to determine the extent of Fe-bkng

intermetallic formation, ouuide of laboratory researchll.

CHAITER ONE : INTRODUCTION 3

1.4 Objective of this Thesis Project

The objective of this thesis project is to show that a method of quantifj4ng the formation

of GFeSi phase, using the thermal analysis methoci is possible. This method will involve

the use of time based and temperature bascd parameters associated with the A&FeSi

thermal effect. The effects of the formation of A&FeSi on other thermal events, such as

the aluminum-silicon eutectic plateau, will also be investigated. Success in this project

will provide an alternative method for measuring iron concentration in alurninurn foundry

alloys.

CI-MTER TWO : IRON IN ALUMINUMSIUCON FOUNDRY ALLOYS 4

Chapter Two

Literature Review of Iron in

Aluminum-Silicon Al10 y s

- --

2.0 Sources of bon in Alurninum Alloys 2.0.1 High Pressure Die Casting

The aluminum die casting industry has been forecast to grow at a rate of 1.6 96

per year over the next few years with a peak of 1 million tons in 199812. Currently

aluminum castings are found in a large number of products ranging from sporting goods,

children's toys, household appliances, office equipment and automotive components. The

duminium die casting industry accounts for up to 113 of aU metai castings made and 60%

of al1 aluminum castings made". The positive attributes of the die cast process are the

volume of castings which cm be achieved in a shoa time (30 shotdrnin), excellent

dimensionai repeatability, ability to cast thin waiis, and excellent surface finish".

Iron in aluminum cast aiioys may be added as an alioy addition by the ingot

manufacturer if the aiioy is intended to be used for high pressure die casting. It is

believed that iron deters the so lde~g of the cast part to the die, but there is some

wntroversy over whether this is actuaUy the case. When a casting and die adhere to one

another via soldering it can be very costly to the foundry, and any measure to deter this

from happening will be advantagrnus. For example, a single cavity die for a die casting

unit with no sides and simple geometry wiiî cost at a minium $50,ûû@3. It has been

reported that a new tool steel reinforced with tungsten carbide particles (composite) cm

provide a higher resistance to attack by molten aluminum, making it ideal for die casting

dies1'. Thus the need to deiiberately add iron to the die cast aüoy may not be as

necessary in the fiiture.

Ejector - die

Die cavity

Cover

Mal ten aiwninum

. Pot

Figure 2.0a : High pressure die castiag unit using goostncck rncthod The goosaicck portion of the unit, made of steel. is ccmtinuaiiy croded away by the aluminum m e l ~ The melt evmtually becomes contamùiated with iron, thus degradhg the casting.

Moving die Fixed die

* - Elements 0 I I a

Air or gas inlet Stal k

ten aluminum

Figure 2.0b : The low-pressure casting unit. The dalk U commoniy made of cast imn, dowing contamination of iron in the moltai metal in the pot

CHAPTER TWO : IRON IN ALUMINUMSILICON FOUNDRY ALLOYS 6

Another prominent source of iron in aluminurn-siliwn die castings is iron pickup

by the aluminum melt h m the pot in which the molten meial is contained. A schematic

of the hot chamber die casting unit is shown in figure 2.h3. It is known that the molten

dUXIilI1um will dissolve iron h m the gooseneck portion of the unit, and graduaily

contaminate the aluminwn melt within the pot. Currently this die casting method is

bccoming more infnquently used h p l y because degradation of the gwseneck requires

its replacement after a certain period of time and reduces the quality of the mele.

2.0.2 Low Pressure Permanent Mold Casting

Low pressure permanent mold casting is a relatively new casting technique which

produces high quality castings. A schematic of such a casting unit is shown in figure

2.0b3. W1th the mold cavity closed, low pressure air (between 2 and 15 psi) is admitted

into the furnace and acts on the surface of the molten metal in the crucible. The molten

metal moves vertically from the crucible, through a cylindrical stalk made of cast iron,

and then entus the mold cavity. The molten metal then flows into the mold cavity slowly

from the bottom allowing for little air to be trapped. When the mold is totally fiîled the

pressure is released and the remaining still-molten metal in the stak f d s back into the

melt. The benefit of this process is that the casting needs no nser, thus increasing yield.

The contamination of the melt by iron occurs as molten metal in the stalk runs back and

mixes with the melt in the pot. Thus, over time, the melt becornes gradually more

contaminated, resulting in the last casting made from a melt having a higher iron content

than the first one made. Using ceramic stalks will remedy this problem, but such stalks

are expensive.

2.0.3 Altuninum Recyding

Another major source of iron contamination, which may becorne more prominent

in the fi-, arises from recycled scrap. Because secondary aluminum has been remelted

several times the potentiai for iron pickup over time is great. In the 14.6 miilion cars

produced in 1994 an average of 72 Ibs of secondary aluminum was usedl3. The amount

of secondary ingot used in foundries is expected to grow drarna t idy in the future due

to the lower cost and increased supply.

CHAPTER TWO : XRON IN ALüMINüMSlLTCON FOüNDRY ALtOYS 7

2.1 Iron-bearing intermetallics in AluminumSilicon alloys

Iron is the most cornmon impurity element found in most aluminum-silicon cast

all~ys~**~~*~. It is a naturai impurity in aluminum ore M e s (bauxite), but can becorne

more wntaminated in foundry processing. When aluminium is in the liquid state (i.e the

casting process) the solubüity of iron is much greater than in the solid state (= 0.052%

by weight at 66û"C)16*17. As a consequence, Fe forms intermetallic compounds with Al

and Si, when the aiioy soiidifiw.

Usually the effect of excess iron content in aluminum-silicon aiioys is the

formation of the thermodynamically stable &phase, having a stoichiornetry of Al,FeSi,

shown in the pertinent part of the Al-Fe-Si temary phase diagram in figure 2. la17. The

pseudo-binary AI-Fe phase diagram, having 8% wt. constant siiicon, is shown in figure

2. lbi7. A15FeSi (25.6% Fe, 12.8 % Si), has the accepted probable ranges of 25-30 96 Fe,

12-15 96 Si, and is monoclinic with lattice parameters a=b=6. 12x1(Tnm, c=41 .5~10~~m, and , ,g 10 18.19.20.21.2223

At smali undercoolings, the growth of this phase evolves via the twin plane re-

entrant edge (TPRE) mechanism (see figure 2.2)". Twin planes in solidifjmg AisFeSi

are aligned parailel to the growth direction. Their intersection with the solid-liquid

interface is the point at which two dimensional nucleation and growth occurs.

Crystailization behaviour such as this usually ressicts the ability of the crystal to bend

or twist during growth, resulting in these interrnetallics having a platelet morphology

which is seen as needle like in 2-D metallographic cross-section1e2. Also, because of the

highly faceted nature of AbFeSi, significant strain in the aluminum matrix just adjacent

to ihis phase exists. As a consequence, under tension the A1,FeSi precipitate may

decohese from the a-aluminum ma&.

The formation of other Fe-rich interrnetaiiics is possible, but depends on the

cooling rate and impurity content of the alIoy7. One which has ken studied extensively

is the a-AI12F~SirA.i,FeSi,Fi2 intermetaUic4. A18FqSi (3 1.6 96 Fe, 7.8 96 Si) and Al12F~Si2

(30.7% Fe, 10.2 % Si) have the probable composition ranges of 30-33 % Fe, 6-12 % Si.

The phase is hexagonal with lattiw parameters of a= 12.3xl@"h, c =Z6.3xl@10m18*1920-n.

CHAPTER IWO : IRON IN AL~MINUMS~LICON FOUNDRY _er-~o YS 8

Figure 2. l a": Liquidus Surface

Al-Fe-Si System

O 1 2 3 4 5 6 7 8 9 1 0 1 1 1 2 1 3 1 4 1 5 Silicon, percent wieght

Figure 2. lb 17: A-Fe Pseudo-Binary Phase Diagram Vertical Section at Constant 8.0% Siliocn

640

620 A

% 600

-

-

'' Al + Liquid - - Al + Si + Liquid

540

520

- Al +P + Si

-

500 I I 1 I

O 0.5 I 1.5 2 2.5 Iron, percent wieght


Twin Plane Re-entrant Edge (TPRE) Growth Mechanism

Monoclinic crystal structure c axis = 20.8 Angstroms a & b axis = 6.12 Angstroms

Atorns attach on the solidlliquid interface in a two dimensional fashion

growth direction

Figure ~ 2 ~ ~ : In this growth mode for Al FeSi intennetallic, twin planes are aligned paralie1 to the growth direction. The intersection of the solidAiquid interface provides sites for 2-D nucleation and growth.


This interrnetallic usually appears as a chinese script in the rnicrostni~ture~~. The a-phase

exerts a l a s deleterious effect on the physical properties of the cast part due to a more

compact shape and a more diffuse interface with the aluminum matrix, resulting in better

cohesion? It has k e n found that this intennetallic fonns either naturally in the later

stages of aiioy solidification when the remaining liquid is depleted in silicon, or in the

presence of impurities such as Mn, Cr, Co, Be and ~ o ~ ~ - ~ ~ . The less common iron-

bearing intermetallics found in aluminum-silicon casting alloys are &aFeSi: and P-

A I , M ~ , F ~ S ~ ~ ~ * ~ . The conditions under which they form are not very clear at present.

2.1.1 Effect of fl(AlFeSi) Phase on the Properties of AlSi Alloys

2.1.1 .O General Properties

The morphology of the 8-phase allows it to act as a stress raiser, consequentiy

underminhg the mechanical properties of the cast part. The threshold amount of iron,

leading to the formation of pnmary MPeSi, that can undermine the properties is > 0.7 % wt. When the phase forms in the eutectic structure (Fe < 0.7 % w) it is believed to

even slightly enhance tensile propertiesZ6. However it must be noted that the percentage

of iron quoted to form primary or secondary AhFeSi depends on cooling rate and silicon

content. The effect of increasing Fe is to gradually reduce the elongation, impact strength

and tensile strength of alurninum-silicon a l l ~ y ? ~ . ~ , while it has been reported that Brinell

hardness and yield strength gradually increase2! The AisFeSi phase has also been

reported to impede fluidity and feeding, and to promote shrinkage por~sityll*~~. The

resulting porosity in itself can also contribute to the deterioration of the mechanical

properties. Fatigue strength can be affectecl by the formation of AisFeSi because the

cohesion between the duminum matrix and the intermetallic in question is poo~?~.

2.1.1.1 Effect of Fe-intermetallics on Corrosion

The effect of Fe on the corrosion properties of AbSi doys has not been weii studied;

however, there are poorly documented reports to indicate that corrosion resistance does

decrease with increasing Fe content for an AL8 %Si-3 %Cu alloyl'.

NOTE TO USERS

Page(s) not included in the original manuscript and are unavailable from the author or university. The manuscript

was microfilmed as received.

PAGES

This reproduction is the best copy available.

2WO : IRON TN ALUMINUMSILICON FOUNDRY -YS

2.1.1.6 Effed of Fe-intermetallics on Elevated Temperature Propertfes

One study for an M i piston aUoy fomd that the m m temperature tende strnigth mis

reduced by the p m c e of the AlPeSi phase, however at elevated temperatures the

presence of these phases increased tensile strength".

2.2.2 Methodology for Correct h g the Deleterious Effects of UsFeSi 2.23.0 Mister Aiioy Addition

Manganese has very litîie solubility in aluminum, and as a consequence Mn in the

remaining liquid will combine with Fe, Si and Al to produce a different Fe-bearing

intermetallic having the stoichiometry Al,,(~n,Fe),Si,'~-". The morphology is the so-

called chinese script, but the phase has been shown to have other morphologies such as

bloclq and dendntid2. The commonly accepted ratio of Mn to Fe to adequately convert

8-phase to a-phase is 1 to 211efi3. Overail the a-phase, as it is commonly calleci, is much

more compact and less detrimental to the mechanical properties and slightly improves

feeding into interdendritic channels. The reason for this formation, in the presence of

Mn, can be understood by the phase diagram of figure 2.3'.

More recently it has been found that strontium, wmmonly used as a eutectic

modifier, has b e n linked to the formation of another a-phase having the stoichiometry

~ l ~ F q S i ~ ~ * ' ~ * ? Apparently Sr reacts with P, preventing heterogenous nucleation of

A1,FeSi on AS a conseyence Al,F@i is formed during eutectic growth when the

Si in the remaining Iiquid is nearly depleted. The method accepted now is to partially

recover the physical properties of aluminum alloys containing Fe which could crystallize

the &phase form via addition of a master aUoy containhg Mn (Le Al-25%Mn, Al-

10 %Mn2 % T ~ Y ~ or Sr. Other foundry researchen have found that master alloys

containing Cr, Co, Be or Mo can have the same effe~t''*'~ of wnverting @-phase to a-

phase.

Figure 2.3 : Simplifïed phase diagram of the ALFe-Si-Mn system with constant M n levels!

2 Mn contnet = 0.2%~~

Alurninum

Silicon \ 2

M n contnet = 0 . 4 % ~

cm)

wt% si1 icon

CHAPTER TWO : IRON Di ALLMINLMSILICON FOüNDRY ALLOYS 14

2.2.2.1 Liquid Alloy Superheat

An alternative method for achieving the predominance of a-phase over @-phase is to

increase the superkat h m 730°C to 9000C31P or greater for most 300 series ailoys.

However this is never done in practice since at these elevated superheat temperatures

serious hydrogen pick-up and oxidation occun. The reason for this phenornenon is that

the heterogeneous nucleant for A l w i is y-N2O3. This oxide changes crystal smcture

at around 905°C to a-A1,4, which is not able to heterogeneously nucleate the A1,FeSi

intennetallic phase effectivelf3. As a consequence A1,FqSi forms instead during the later

stages of eutectic fkezing.

2.2.2.2 Cooling Rate

Cooling rates have dso been shown to have an effect on the form of Fe-bearing

intermetallic. When Mn contents are minimdly small, the Al,FeSi phases are affectai

by cooling rate alone. The nucleation temperature of A&FeSi is suppressed with

increasing cooling rate for a given iron content. Thus the time available for the QFeSi

phase to grow is shortened and the overall length of the QFeSi needles themselves is

reduced.

The formation of A1,FeSi at faster cwling rates when rnanganese is present is

slightly different. This formation of AlPeSi phases in the presence of Mn may be

descnbed as foilows. According to the Al-Fe-Si-Mn phase diagram (figure 2.3~)~-''

solidification begins with the growth of dendrites, labeiîed as arrow 1. The remaining

liquid becornes ncher in solute elements (Fe, Mn & Si), then the segregation line (arrow

2) penetrates the AlI,(Mn,Fe),Si2 field, where A1,5(Mn,Fe)3Si2 particles grow within the

liquid. Then at the dashed region ( m w 3) both Ai15(Mn,Fe),Si2 and A1,FeSi phases

grow together. Eventuaîiy AbFeSi grows with elemental Al and Si at the temary eutectic

(point 4).

The relative amounts of a(AlFeSi) and P(All;eSi) that are fomed in the

microstructure depend on the cooiing rate and Fe/Mn ratio. At slower cooling rates Fe

would be consumed by the formation of A1,,(Fe,Mn),Si2, while at fast cooling rates not

al1 of the Fe is absoroed in passing dong arrow two. Thus some AiPeSi will form. As


a result there is a fraction of Fe to be absorbed in solidification in the Ai,(Fe,Mn)&

phase field, and then d u ~ g the later stages of solidification, the formation of G F e S i

becornes favourable.

2.2.2.3 Solution Heat Treatment

Recent research has discovered that solution treatment of a 356 dioy for 12 hours at

515°C foiiowed by heating at 540°C for a further 12 hours resulted in the aimost

complete dissolution of ~ l ~ ~ e S i ? Tensile and hcture testing was done and improved

properties were found. Similar results were found by Croweii and Shivkumar", except

that the chinese script, uniïke the needle phase, underwent no change during solution

treatment. It should be pointeci out that both studies were limiteci to A&FeSi that was

secondary , not primary . No real study has been done on alloys that form pnmary A1,FeSi

since these alloys are usuaUy designated for die casting, and these casting types are

usuaiiy never heat treated.

2.2.2.4 Dilution

Primary or secondary ingot producers do have the option of adding enough aluminum,

siiicon, magnesium and copper to a melt so that the ove& iron contents are reduced

simply by the dilution effect.

CHAPTER THREE : THERMAL ANALYSE OF ALUMlNüMSILICON ALLOYS 14

Chapter Three

Thermal Analysis of

Alurninum-Silicon Alloys

3.0 Thermal Analysis 3.0.1 Introduction to Basic concepts

In typical thermal analysis, sarnples are pou& into a srnd crucible or cup and solidify

at a specific cooling rate. The temperature is remrded (using a sheathed thermocouple)

with real time from the liquid state, through the solidification range, to the solid state.

The resulting plot (temperature vs. time) is the cooling curve. This cooling curve can

then yield information about the metallurgical aspects of the sample, and consequently

infer on the possible rnicrostmcture of the casting or ingot king made.

Thermal analysis, originaily used for the cast iron industry, has ben an accepted

method for quality control in the aluminum foundry industry for over ten years. This

methodology is used to gauge the extent of grain refinement and eutectic modification of

aluminum-siliwn aUoys, but it has not been used to quanti@ intennetallic formation (eg.

B(AlFeSi) phases).

When a phase begins to grow in the melt, the cooling rate of the solidiQing ailoy

changes and is seen as slight inflections, or dramatic changes ( Le. increase in

temperature, caiied recalescence) on the cooIing curve, in what would othenuise be a

smooth curve of wnstantiy decreasing temperature with time. Large volume phases

which evolve during solidification often lead to d e s c e n c e on the oooling curve, while

smaii volume phases may show up as only a thermal anornaly. Certain phases have

undercoolings or heat effects which are very smaii, and consequently are not evident on

the original cooling curve. Then, the use of the first derivative curve can be employed

CHAFïZR THREE : THERMAL ANALYSLS OF ALUMINUMSILTCON ALLOYS 17

to accentuate these heat effects.

3.0.2 T h d Analysis of a Typical Ai-Si PUoy (Example of 319.2 Alloy)

The 319 alloy has been offen studied using thermal analysis to study the formation of

Uon-bearing i . n t e r m e t a l l i c ~ ~ ~ ~ ~ ~ ~ ~ ~ - * . An example of a coolhg curve and its first

denvative curve for 319 aUoy is s h o w in figure 3.0. The first thermal effêct is seen at

about 604"C25, where the coohg rate slows down considerably , then reverses (dT/dt > O) as undercoohg and recaiescenœ occur. This characterises the growth of alurninum

nuclei into dendrites. The magnitude of undercooling and its duration are detennined by

the initiai grain s*. Once the cooling rate again becornes negative (dT/dt < O), the

dendrite structure begins to coarsen. At 565°C another undercooling effect occurs,

foliowed by a long plateau where the cooling rate is almost zero (dT/dt = O)? This

represents the growth of the eutectic structure withui the interdenciritic regions. D u ~ g

the time for eutectic growth the dendrites continue to c o a r ~ e n ~ ~ * ~ ~ . The final reaction in

this alloy is the temary eutectic reaction leading to formation of &Cu at about 520°C.

Aluminum-Silicon Cooling Curve & Derivative Curve 3 19.2 Alloy

1

Eutectic plateau, dT/dt - O

- - Tirne (sec) Figure 3.0": The typical coohg curve and associated derivative curve of an alLuninumsilicon (3 19) doy.

E : THERMAL ANALYSIS OF ALtJMïMMSfLICON ALJDYS 18

3.1 The Thermal Analysis of Minor Reactions 3.1.0 j3(AiF'eSD Reactions

In A356 alloy the Fe amounts are very low and the formation of &FeSi and A&FqSi,

occurs during the end of the eutectic reaction. At that point most of the remaining liquid

is depleted in Si and enrîched in iron, favouring some formation of A&Fe+Si, the so

called a-phase. The heat effect h m this reaction is very smaü and is not distinguishable

h m the eutectic reactionu. As a consequenœ the heat effect is not observable for Fe

contents of about 0.7% wt or less since it becornes submerged within the heat effect of

the eutectic plateau. For iron contents greater than 0.7 1 wt, the formation of Al5FeSi

ieads to a heat effect that is observable before the eukctic formation, but after the

primary arrest as seen in the 3 19 alioy coohg c w e and derivative curve in figure 3.1.

The A15FeSi phase can also form in a posteutectic reaction, as in the case for 3 19 aiioys

where the reaction at 520°C as reported in the l i t e r a w 3 is:

Aluminiun~Silicon Cooling Curve & Derivative Curve 3 19 d o y with high iron (O. 8% wt)

- - T h e (sec) Figure 3. f' : Cooling curve of an 3 19 alloy having a high iron contnet

CHAPTER THREE : THERMAL ANALYSE OF ALUM][NUMSILICON ALLOYS 19

3.1.2 Other Post-eutectic Reactions

The most cummon and observable post-eutectic reaction is the one involving

magnesiumz:

L-> A l + Si +Mg2Si (1.1)

This occm as a anomaious bump after the eutectic plateau at a temperature of about

55PC for 356 doy.

Post eutectic reactions occur in 319 alloy, which has copper contents of about

3.5% by weight. Three possible raidons involving copper can occur, as reported in the

l i t e r a t ~ e ~ ~ . The first is a simple reaction:

L-> Al + &Cu (1.2)

This reaction m u r s at 549°C and is usually observable on cooling curves of alloys that

are not modified. The second one is reaction (1.0), which cm be seen at 520°C. It is

during this reaction that most of the 4 C u forms. The third reaction, h a h g a noticable

heat effect is observed at 500"CY:

3.1.3 Thermal Analysis Parameters

Time and temperature parameters can be measured to gauge grain refinement and eutectic

modification. However the cooling curves used for this cornparison must be of the sarne

melt mass and configuration to ensure that the same cooling rate is attained. Figure 3.1

shows the typical time and temperature parameters of various arrest points that can be

measured in typical alurninum-siliwn aiioys. Cwling rate may be calculated from the

dope of the cooling curve itself between primary and eutectic anest points". For primary

growth, to gauge grain refinement, the undercooling can be measured as the temperature

difference betwcen the minimum of undercooling and the temperature of equilibrium

transformation*. For eutec tic modification the eutectic undermihg temperature

parameter is defined as the temperature difference betweui the minimum temperature of

undercooling and the temperature of the eutectic plateau. In practice most foundries

CHAPTER THREE : THERMAL ANALYSE OF ALUMINUMSTLICON ALLOYS 2Q

Theml Analysis Pararnerters

Liquidus undércoo + 4- Reference cooling / !-

Liquidus undercooling ("C

tirne (sec) Fig 322528 : A coolhg m e of a 356 alioy and the most commonly rnePsured parameters .

Thermal Analysis Time Paramerters

( Total solidification time -1 \

100 200 300 time (sec)

Fig 3.338 : A coohg nwe of a 356 dioy and timc parameters proposeci by m e r s as given in the literature.

simply measure the depression of the eutectic plateau, AT, which is defined as the

temperature difference between the eutectic plateau arrest temperature for the base alloy

and that of the same aUoy in modified condition?

Time parameters, as show in figure 3.2, can only be teliably measured if the

cooling rates of aî i thermal analysis samples are the same. The use of time parameters

has been investigated for eutectic modification of 319,355,356,357 and 380 alloys*?

It has k n found that with incfeasing strontium content the time of the eutectic plateau

increases. The start and end times were determined by inflection points associated with

the latent heat of fusion given off from the aluminum-silicon eutectic growth. A generai

time parameter for the whole cooiing curve for Ai-Si aiioys was suggested by GO&'

and Spark~nan~~, but was never measured quantitatively as part of their research. The

examples for suggested time parameten are seen in both figures 3.2 and 3.3.

The use of derivative curves higher than first has only been report& for cast iron

where the first and second denvative is used to gauge the formation of graphite phases,

predict graphite or carbide eutectic, ratio the percentage of austenite venus percentage

of eutectic, predict mottle in rnaiieable iron and predict carbon quivalent. Figure 3.4

shows a near eutectic ductile iron cooiing curve with the associated fxst and second

denvative curves4'. A point of interest is that for cast iron the beginning of eutectic

formation is signified by the maximum of the second derivative and is labelled in figure

3.4 as MXSDES. Another point of interest is the minimum in the second derivative

associated with the end of eutectic fieezing, iabeiied MXCREE. These uses of the second

denvative curve for cast iron, although not commonly measured on the foundry floor,

wiii be of use in the thermal anaiysis of Al-Si ailoys in this thesis. To date only the first

denvative has been used to interpret AI-Si cooling curves, and no parameters other than

cooling rate or accentuation of difficult to detect thermal events has been inferred from hem24.3337

One report however proposed interpretation of the second and third derivative to

signal the presence of AifieSi phase for 319 ailoy. The cooiing curve in question had an

A1,FeSi arrest point not detectable on the original cooling curve, but identified by the

zero crossover in the third derivative and minimum of the second derivative in tandem39.

CHAPTER THREE : THERMAL ANALYSE OF ALUMINUMSILICON ALLOYS 22

Reaüstically the zero crossovers of higher denvative curves are numerous and may result

in erroneous interpretation. Also this method may be useful for identifying A1,FeSi if i t

occurs as a difficult to see arrest.

Figure 3.4 : Typical thermal analysis curves of a near eutectic ductile cast iron with asmciarrd first and second

derivative curves*.

3.1.4 Limitations of Thermal Analysis

While thermal analysis can be a powerful tool for gauging the potential microstructure

of a casting some serious limitations do exist. Thermal anaiysis is usudly a batch

process. For example, to accurately rneasure the suppression of the eutectic plateau due

to modification the temperature of the unmodified eutectic plateau of the same melt must

be determined. Using the base unmodified temperature from a different melt can

introduce an erroneous determination of modification because of slight chemistry changes

(eg. phosphorus content). The change in eutectic temperature is typically SOC to 8°C.

while thermocouple accuracy is approximately within & 2OC. Finally , thermal anal ysis

can be used usually at only one cooling rate, while casting or ingot microstructure is

determined by a variety of mling rates dependhg on casting or ingot shape and complexity.

Chapter Four

Experimental Methodology

4.0 Chemistry & Microstructure of AlSi Alloys Studied The main alurninum-silicon alloys studied in this thesis were 319, 356 and 413. For each

aüoy type the chernistry and microstructure will be discussed to show the importance of

Fe content and the potentiai for AlPeSi phase formation. The Aluminum Association

typicaliy designates the higher puriq (Le lower Fe contents) alloys with a prefix letter

before the alloy identification number (eg. A356 alioy). The primary alloying element

in ail of these alloys is siiicon. Silicon gives aiuminum in the liquid state the fluidity

needed for good castability and helps to reduce solidification shnnkage of the overall

casting. Other alloying elements will be discussed in their appropriate sections, as

needed.

4.0.1 319 (Al-Si-Cu) alloy

The major ailoying elements of this alloy type are silicon, mangrnese, copper and iron.

Copper strengthens the alloy through solid solution strengthening of the aluminum

matrix. Complex intermetallics also fonn in pst-eutectic mictions during the final stages

of solidification. The iron content varies depending on the intended use for the alloy. If

the aUoy is to be used in die casting, a grade of 319 having a high level of iron will be

used, while for structural castings a higher purity grade of 3 19 is employed. Manganese

content is intended to prornote the formation of Al,,(Fe,Mn),S&. This more desirable

morphology becornes possible when the Fe:Mn ratio is 2: l 1 l * I 5 . The o v e d chernistry as

defined by the Alurninum Association is show in table 4.0". As can be seen from table

4.0 the major differences between the different grades of 319 ailoy are the iron,

CHAPïER FOUR : EXPERIMENTAL METHODOLOGY 24

mangrnese and zinc contents. Notice that for the higher iron containhg 319 aiioys the

manganese content also increases. The microstructure of a t y p i d 319 aiioy is shown in

figure 4.h. The mostiy giey background represents the dendrites of aluminum. The dark

acicuiar phase is sificon which forms through the eutectic reaction, whiie the lighter grey

needle Iike-phase is A1,FeSi. This intermetallic occua as prirnary (formed prior to

eutectic formation) and secondary A15FeSi which forrns as part of an Al-Si-Fe eutectic

reaction or within the main Ai-Si eutectic. The final phase to form is the CUAI,

intermetallic which may have a blocky morphology or be present as finely dispersed Al-

CUAI, eutectic.

Table 4.0: Aluminum Association Specifications for 3 19 Alloy.

II I Composition, wt. pct.

4.0.2 (Al-Si-Mg) 356 alloy

The pnmary alloying elements in this ailoy are silicon and magnesium. The addition of

0.3% magnesium allows for heat treatability of the ailoy, improving mechanical

properties. The 356 aUoy is a very important structural aUoy usuaUy having a lower iron

content which increases its tensile and impact strength as weii as its elongation. The 357

alloy is very similar to 356 except that the magnesium contents are slightiy higher

( 0 . 5 5 ~ 46). The typical chernical composition of the various grades of 356 aiioy as

ÇHAPTER FOUR : EXPERIMENTAL METHODOLOGY 25

outlined by the Aluminum Association are given in table 4.141. The main difierences

behveen these grades are in the iron and rnagnesium content.

The microstructure of 356 alloy is show in figure 4.0b. This microstructure is

simihr to that of 319 aiioy, having a mainly grey background representing the aluminum

dendrites and the acicular silicon of the eutectic. The main difference is in the absence

of the CuAl, phase since no copper W s t s as an alloying element for this particular alloy.

Instead some Mg,Si may be observable. The A1,FeSi intennetallic is less prevalent within

the microstructure compared to 319 a o y since the iron content is lower.

Table 4.1: Aluminum Association Specification for 356 Alloy.

Composition, wt. pet.

4.0.3 ( A M ) 413 aiioy

This dloy is a eutectic alloy, and because of the higher silicon content, castings made

of this alloy are usually difficult to machine. The increased siIicon content allows for less

solidification shrinkage, while the latent heat of fusion of siiiwn keeps the melt very

fluid. Thus castings having thin walls or intricate designs can be made using this dey.

There are no other alloying elements used except for iron which is sometimes added by

the ingot producer if the alioy is intended for use in die casting. The iron contents of the

4 13 aUoy is usually quite high, particularly for 4 13.0 ailoy , and the formation of the Fe-

bearing intemetallics is quite prorninent. The typical chernical compositions of the

different grades of 413 alloy as outlined by the Aiuminum ASsociation are aven in table

4.241.

The typical microstructure of 413 is seen in figure 4.0~. In the non-modified state

an acicular eutectic silicon is observed, weii dispersed within the aluminum rnatrix.

Because of non-equilibrium cooling conditions and slight chemistry changes, polyhedrai

Silicon (or primary silicon) can form in the microstmcture. The effect of a modifier can

result in making the alloy solidify in a slightly hypoeutectic manner leading to the

presence of primary aluminum dendrites. Long grey needle-like A15FeSi intennetallic are

observable within the microstructure of figure 4.0~.

Table 4.2: Aluminum Association Specifications for 4 13 Alloy .

MOY

413.0

413.1

413.2

A413.0 - A413.1

A413.2

B413.0

Composition, wt. pct. l

Si

11.0-13.0

11 .O-13.0

11 .O-13.0

11.0-13.0

11 .O-13.0

11 .O-13.0

11 -0-13.0

Fe

2.0

0.7-1.1

1.3

1 .O

0.6

0.5

0.4

Cu

1 .O

O. 10

1 .O

1 .O

O. 10

O. 10

O, 10

Zn

0.50

O. 10

0.50

0.50

0.05

0.05

0.05

Mn

0.35

0.10

0.35

0.35

0.35

0.35

0.35

Mg

O. 10

0.07

O. 10

O. 10

0.05

0.05

0.05

CHAPTER FOUR : EXPERIMENTAL METHODOLOGY 27

Figure 4.0a : Microstructure of 3 19 ailoy. I .) a-aluminum, 2.) silicon, 3.) AI,Cu-Al-Si

eutec tic, & 4. ) A1,FeSi in termetallic (needle scnp t)

Figure 4.0b : Microstructure of 356 alloy. 1 .) a-aluminum, 2.) silicon & 3.) A1,FeSi

(needle script)


Figure 4 . 0 ~ : Microstructure of the 4 13 alloy. 1 .) a-alurninum, 2.) silicon, 3 .) A15FeSi

(needle script) & 4.) Al,,(Mn, Fe),S i, (chinese script)


4.1 Experirnental Foundry Procedure To exploit the use of thermal analysis as a tool for quantifying iron content, various

deys were studied via th& analysis while increasing the iron content from the base

ingot amount to the highest iimits dehed by the Aluminum Association. A 10 Kg mass

of cut AbSi alloy ingot was melted in an dumina-silica (%l%&03/10%Si03 mucible

using an electric &stance fumace. Once the alloy was totally molten the temperature

was maintained at approximately 730°C. Before any master alloy additions were made

to the melt, three thermal analysis samples were poured to obtain cooling curves

representative of the base alloy chemistry. Master dloy additions, via an Al-24.796Fe

master aiioy were then made to increase the Fe content of the main melt by

approxirnately 0.15% weight increments. Fifieen minutes was ailowed for dissolution

while stirring was performed during 113 to 112 of this time. Three more thermal analysis

samples were then poured for each of the composition increments. This process of adding

master dloy and plotting of three cooling c u n e s after each melt addition was done up

to an Fe content equal to the maximum Iimit defined by the Aluminum Association.

Finally it should be pointed out that after each aUoy addition spectrochemicai samples

were poured so that the actual chemistry of the melt could be determined.

4.2 Experimental Setup for Thermal Analysis The primary components needed to do basic thermal analysis are a personal cornputer,

12 or 16 bit analog/digital interface, and a crucible of melt with a shathed thermocouple

(K type) immersed into the melt. A schematic of this setup is shown in figure 4.1. For

the majority of samples a stainless steel crucible having about 1 inch of fiberfrax

insulation around the sides and bottom was used as show in figure 4.2a. This thermal

analysis sample sehip was designed to achieve a cooling rate of 0.1O"Clsec over the

region between initial primary a-aluminum formation and the beginning of the Al-Si

eutectic. This wiil be referred to as the insulated cup method. For sarnples which had a

reference cwling rate of 0.4S0C/sec, the thermal analysis samples were poured in silica

sand (figure 4.2b). An open cavity in the silica sand was formed using an uncut thermal

Figure 4.1 : Thermal Amlysis Setup.

Figure 4.2 a): Stainless steel crucible with fiberfrax insulation to achieve the O. 1 O°C/sec referenec cooling rate. b.) Sand sand mold to achieve 0.4S°C/sec.

ThtImmuple - (stainless steel sheath on end) -

CHAPTER FOUR : EXPERIMENTAL METHODOtOGY 31

anaiysis sample from the steel insulated cup used for the 0.1O"Clsec reference cooling

rate. This method will be referred to as the sand mold method.

The software used for data acquisition was the MeltlabTM software developed by

Foundry Information Systems. This software worked in conjunction with a 16 bit

anaiog/digitd converter. This software dso produced the cooling c w e s and derivative

curves used in this thesis. The analysis of the coolhg curves was done using the

G9ûûûN software also developed by Foundry Information Systems. The G 9 0 m

software is a research onented software allowing for zoom ups of various portions of the

cooling curve and its possible five derivatives. One of the main reasons for using this

particular software is that cooling curve data and derivative curves could be imprted

from the Meltlab software and be converted into an ASCII file. Once the data was in

ASCII format it could be imported into Freelancem for Windows for graphitai

presen tation.

4.3 Post Foundry Experimental Analysis

It is important to correlate aspects of the AbFeSi thermal analysis from cmling curves

with other methods for quantifying Fe content. For this correlation, image analysis and

spectrochemical anal y sis were used .

4.3.1 Spectrochemical Analysis

To veriQ the chemistry of the alloy in question with iron additions, spectrochemical

analysis was performed on 70 gram disc-like aluminurn alioy samples using a Baird-

A tomicm spectrome ter Mode1 DV-2. Be fore performing the spectrochemical anal y sis

each solidified 70 gram disc was ground using 60, 120, 240 and then 400 grit paper.

Because of the caiibration of the emission spectrometer, iron levels higher than

1.2% wt were not possible to measure. As a consequence, determination of iron levels

above 1.2% wt was perfonned using Atomic Absorption Spectrophotometry.

CHAP'ïER FOUR : EXPERDMENTAL METHODOLOGY 32

4.3.2.1 Poüshing of Aluminum Ailoy Samples

AU thermal analysis samples were cut near the vicinity of the sheathed thermocouple tip.

It is at this region where the cooling curve best reflects the microstmcture. The srnail cut

samples were mounted in cold mount resin and then wet ground with 60, 180, 320, 600

grit silicun carbide paper on a Leco automatic polisher. Fine grinding was done using a

5 and then 0.3 micron alumina (%O3) suspension in water. Prior to fine grinding, a i i

samples were placed in an ultrasonic bath for about 5 minutes to physically remove any

of the alumina that may remain on the surface. Final polishing was then done using

colloidal silica solution. Shrinkage porosity that formed upon solidification resulted in the

entrapment dunng polishing of colloidal silica within the cavities of exposed shrinkage

pores. This trapped colloidal silica would subsequently smear over the surface obstructing

a clear view of the microstnicture under a light microscope. This problem was remedied

by placing all samples, after a colloidal silica polish, into a beaker of methanol placed

in an ulhasonic bath for approximately five minutes. Such cleaning would vibrate any

remaining colloidal silica out of the pore cavities.

4.3.2.2 Image Anaiysis Procedure

Image analysis was performed on the prepared metallographic samples. The objective of

this was to quanti@ the percent area of Al,FeSi present in the microstructure. This was

done using a Leco image analyzer interfaced with an optical microscope through the 2005

image analysis software. This system is capable of distinguishing 256 grey levels. A

histogram was determined, where A15FeSi had a grey shading different from the other

phases within the microstnicture. Each frame was at 2WX magnification and

approximately 10 to 15 frames were taken per sample. A mean value of percent area

A1,FeSi with a standard deviation was then deterrnined from the data.

For the 319 ailoy, the copper interrnetallic and A15FeSi phases within the

microstrucnire were asentially the same wlour, and hence difficult to separate. This

problem was remedied by etching the samples using a solution of 25 ml nitric acid

@NO3) and 75 mL water. This etchant resulted in the copper intermetallic a p p e a ~ g

dark grey and thus clearly distinguishable from the QFeSi phase.

FWE : RESULTS OF THERMAL ANALYSE 33

Chapter Five

Results of Thermal Analysis

5.0 Spectrochernical Analysis of AluminumSilicon Alloys Spectrochernical analysis was performed to determine the exact chernical composition of

the cast alloys used in this thesis beforr iron addition to the melt via an Al-24.71Fe

commercial master aüoy. This is important since other elements such as manganese cm have an effect on the solidification characteristics of the iron-bearing phases that form.

After iron additions were made, spectrochemical samples were also poured so that the

iron content could be correlated to the A1,FeSi phase in the microstructure and its

associated heat effect on the thermal analysis curves. Other elements that are of particular

interest are the siliwn and copper contents since they can have an effect on the

temperatures at which some reactions do occur. Table 5.1 shows the chernical

compositions of the as-received commercial ingots used for the experiments.

Table 5.0 : Chemical Composition of As-Received Aluminum-Silicon Aiioys.

Il I Composition, W. pct.

CHAPTER FIVE : RESULTS OF THERMAL ANALYSIS 34

5.1 Results of Thermal Analysis & Associated Microstmcture

The cooling curves with second derivative and associated micrographs for the three

different aiuminum-Silicon doys studied are pfesented in the following sections. In

pezforming the experiments demibed earlier, it was recugnized that coohg rate was an

important variable affecthg the cooling curve. The cooling rate was determined as the

dope on the cooling curve between the primary thennal arrest and the eutectic plateau.

This value is r e f d to as the reference cooling rate. To explore the importance of

reference cooling rate, the A356 alloy was studied with thermal anaiysis samples having

two reference coohg rates : O. lO"C/sec (insuiated cup) and 0.45OClsec (sand mold).

Repeatability was studied using the 319.2 aiioy on which the cooling curves with

in&g iron contents were done twice. For the third alioy, A413, thermal analysis

samples were poured in the insulateci cup show in figure 4.2a. A series of cooIing

curves for 413 aiioy having a graddy increasing iron content was done once.

The second denvative curve was used rather than the first derivative and is shown

in the cooling curve figures of this chapter. The first derivative is a measure of the

instantaneous cooling rate dong the cooling curve. As a consequenœ the first derivative

has been used in research to signal the presence of nearly invisMe arrests of minor

reactions on coohg curves. The second denvative gives the change in cooling rate, and

can result in peaks when dramatic changes in cooling rate mur. For example, when a

phase begins to grow, the coohg rate changes slightly yielding a maximum or minimum

in the second derivative. The maximum of the second derivative wiII be used as an

important guiding tool in this thesis, signalling the beginning and end of phase formation.

The local maxima of the dq/de c u v e s that are of interest in this thesis are marked by

arrows on the appropriate figures.

5.1.1 356 AUoy (0.10 OC/sec & O.4S0C/sec)

The thennal analysis of A356 aüoy, having a reference cooling rate of O. lO"C/sec, was

perforrned using six different iron contents. For each iron content three cooling curves

were done to check repeatability, leading to a total of eighteen coohg curves. Three

cooling c w e s representing this aiioy, each having three different iron contents are


shown in figures 5. la, 5.23 and 5.3a, and their associated rnicrographs are given in

figures 5. lb, 5.2b and 5.3b. In figure 5. la a cooling curve for A356 aüoy h a h g an

artificialiy increased iron content of 0 . 5 4 % ~ is displayed. This cooling curve represents

the minimum iron content that was detectable for this senes. The basic trend for ail 356

cooling c w e s using this reference cooling rate may be described as foiiows. The coolhg

curve shows its first themial arrest point at 612°C where the growth and formation of

alurninum nuclei m u r . After this arrest, the temperature of the solidifying d o y

continues to decrease. D u ~ g this time the liquid is progressively enriched in silicon. For

the aiIoys shown in these figures, a thermal anomaly occurs at 572°C (0.54% wt Fe) in

figux S.la, 594°C (0.94% wt Fe) in figure 5.2a and 605°C (1.25% wt Fe) in figure

5.3a. This heat effect is caused by the latent heat of fusion of the A1,FeSi phase. The

formation temperature of this phase is clearly increasing since the increase of iron has

adjusted the temperature at which AsFeSi forms. At 577T, for figures 5. la, 5.2a and

5.3a, the cooling rate changes again to nearly zero. This continues for some duration as

the latent heat of fusion for the aluminum-silicon eutectic is evolved. The final arrest

point occm at 550"C, when the Mgsi-A1Si ternary eutectic phase is forrned. The

microstructures for al1 of the associated cooling curves are shown in figures 5.1 b, 5.2b

and 5.3b. The light grey background is the a-aluminum while the dark grey phase is the

silicm of the eutectic structure. The A1,FeSi is seen as a grey needle-like phase in the

microstructure. As the iron content increases, the length of the needle phase increases

due to the increased time available for growth as the formation temperature rises.

Figures 5.4a, 5.5a and 5.6a show three cooling curves of the A356 aiioy each

having different iron contents and solidified at 0.45"Clsec. As seen in figure 5.4a. the

A1,FeSi becomes observable only at 0.75% iron indicating that the limit of detectabiiity

has risen compared to sarnples solidified at the slower cooling xate of O. 1CPClsec. The

formation temperature of this intermetallic increases to 5780C (0.996 wt Fe) in figure

5.5a, while in figure 5.6a the formation temperature has increaKd to 6000C (1.4% wt

Fe). The associated micrographs shown in figures 5.4b, 5.5b and S.& indicate that the

A15FeSi platelets are considerably shorter than those obsewed in figures S. lb, 5.2b and

5.3b, due to the faster cooiing rate.

CHAPTER FIVE : RESULTS OF THERMAL ANALYSE 36

Cooling Curve & Associated Second Derivative for A3 56 : Reference cooling rate = O. 1 O' C/sec

I I l i I I -0.02

O 200 400 600 800 1,000 tirne (sec)

Figure S.la: The cooling curve and the associated derivative c w e for a A356 alloy having 0.54% wt iron.

Figure S. lb : The microstnicture of the A356 alloy solidified according to the cooling curve above. 1 .) B-phase, 2.) silicon & 3 .) a-aluminum.

CHAPTER FIVE : RE.SULTS OF THERMAL ANALYSIS 37

Cooling C w e & Associated Second Derivative for A356 alloy : Reference cooling rate = 0.10 ' C/sec

cnn ' , \ ) 1

J V V

O time (sec)

Figure 5.2a : The cooling curve and associated denvative curve for a A356 alloy having 0.98% wt iron.

Figure 5.2b : The microstmcture of the A356 alloy solidified according to c~o l ing curve above. 1 .) &phase, 2.) silicon & 3 .) a-aluminum.

the


Cooling Curve & Associated Second Derivative for A356 alloy : Reference cooling rate = 0. 10°C/sec

cooling curve second derivative v time (sec)

Figure 5.3a : The cooling c w e and associated derivative curve for a A356 alloy having 1.25% wt iron.

Figure 5.3b : The microstruchue of the A356 alloy solidified according to cooling curve above. 1 .) B-phase, 2.) silicon & 3 .) a-alurninum.

the


Cooling Curve & Associated Second Derivative for A356 alloy : Reference cooling rate = 0.4S°C/sec

cooling cuve second derivative h\ time (sec)

Figure 5.4a : The cooling curve and the associated derivative c w e for a 356 alloy having 0.75% wt iron.

Figure 5.4b : The microstructure of the 356 alloy solidified according the cooling curve above. 1 .) p-phase, 2 .) silicon & 3.) a-aluminurn


Cooling Curve & Associated Second Derivative for A356 alloy : Reference cooling rate = 0.45' Ckec

1 O0 200 300 400 500 time (sec)

Figure 5.5a : The cooling curve and the associated derivative curve for a 356 alloy having 0.90% wt iron.

Figure 5.5b : The microstructure of the 356 alloy solidified according the cooling curve above. 1.) P-phase, 2.) silicon & 3.) a-aluminum


Cooling Curve & Associated Second Derivative for A356 alloy : Reference cooling rate = 0.45'Clsec

U ( cooling m e second derivarive 1 Y

time (sec)

Figure 5.6a : The cooling curve and the associated derivative curve for a 356 alloy having 1.4% wt iron.

Figure 5.6b : The microstructure of the 356 alloy solidified according the cooling n w e above. 1 .) P-phase, 2.) silicon & 3.) a-aluminum

CHAFlZR FIVE : RESUÏ.TS OF THERMAL ANAI.YSTS 42

5.1.2 319.2 M o y (O.lO°C/sec)

The thermai analysis of 3 19.2 aiioy, havhg a r e f e ~ c e cooling rate of O. lCPC/sec, was

performed using seven different iron contents for both trials. For each iron content two

or three coolhg c w e s were done to check repeatability. Three cooiing curves

representing this aiïoy in trial one, each having three diffant iron contents are stiown

in figures 5.7a, S.& and 5.9a, and their associateci micrographs can be seen in figures

5.7b, 5.8b and 5.9b. In figure 5.7a a cooling curve for 3 19.2 alloy having the base iron

content of 0.3696wt is shown. The basic trend of the temperaturetirne curve for this

alloy is similar to that of the 356 aUoy except that the primary aluminum growth

occurred at 604"C, the eutectic formed at 567°C. and the post-eutectic reaction is an

WU-Ai-Si temary eutectic reaction which occun at 520°C. In figure 5.8a the cooling

curve of the aüoy having an iron content of 0.90% wt is shown. The t h e d anomaly

associated with the A15FeSi crystallization occurs at 577OC. In the associated rnicrograph,

the Al,FeSi phase is clearly observed in the microstnicture. Both the silicon of the

alurninum-silicon eutectic and the AI,Cu-Al-Si eutectic is seen growing from the Al,FeSi

phase. Figure 5.9a shows the cwling curve for the 3 19.2 alloy having 1.2 % wt iron. The

A1,FeSi thermal anomaly is now even larger and begins to fonn at 590°C. As seen in the

associated micrograph, the microstructure appears to be the same as before except for

the increased length of the AisFeSi phase. Results for the rrpeat (trial two) are seen in

figures 5.10a, 5.1 la and 5.12a, and the microstructures, seen in figures 5. lob. 5.1 1 b and

5.12b. are essentially the same as in trial one. More wiil be discussed about repeatability

in subsequent sections of this thesis.

5.1.3 A413 AUoy

The thermal anaiysis of A413 alloy , using the insulated cup method, was performed with

seven different iron contents. For each iron content, two or thne cooling curves were

taken to check repeatabiiity. Three cooiing curves representing this aiioy, each having

different iron contents are shown in figures 5.13a. 5.14a and 5.15a and their associated

micrographs are seen in figures 5.13b, 5.14b and 5.lSb. The cwling curve for this

eutectic alloy having the base iron content is composed of a single long plateau

coreresponding to aluminum-siliwn eutectic growth at 577°C. The microstnicture reveals


Cooling Curve & Associated Second Derivative for 3 19.2 Alloy (trial one) : Reference cooling rate 0.1 O°C/sec

time (sec) Figure 5.7a : The cooling curve and associated derivative curve for a 3 19.2 alloy having 0.36% wt iron.

Figure 5.7b : The microstructure of the 3 19.2 alloy solidified according to the cooling curve above. 1 .) P-phase, 2 .) silicon, 3 .) Al 2Cu-Al-Si eutectic Br a-aluminum

CHAPTER FXVE : RESULTS OF THERMAL ANALYSTS 44

Cooling Curve & Associated Second Derivative for 3 19.2 Alloy (triai one) : Reference cooling rate = 0. 10°C/sec

-

cooiing curve second derivative

time (sec)

Figure 5.8a : the cooling curve and associated denvative curve foe a 3 19.2 alloy having 0.90% wt iron.

Figure 5.8b : The microstructure of the 3 19.2 alloy solidified according to the cooling cuve above. 1 .) P-phase, 2.) silicon, 3 .) AI, Cu-Al-Si eutectic & a-aluminum

CHAeTER FIVE : RESULTS OF THERMAL ANALYSIS 45

Cooling C u v e & Associated Second derivative for 3 19.2 AIloy (trail one) : Reference cooling rate = 0.1 O°C/sec

F

- - -

t I 1 I I I I I -0.01 O 200 400 600 800 1,000 1,200 1,400

tirne (sec)

-

-

cooling curvt second dcrivativc

Figure 5.9a : The cooling curve and associated derivative curve for a 3 19.2 alloy having 1 2% wt iron.

Figure 5.9b : The microstructure of the 3 19.2 alloy solidified according to the cooling curve above. 1 .) B-phase, 2.) silicon, 3 .) AI, Cu-AI-Si eutectic & a-aluminum

CHAPTER FIVE : WULTS OF THERMAL ANALYSIS 46

Cooling Curve & Associated Second Derivative for 3 19.2 alloy (trial two) : Reference cooling rate = 0.1 O°C/sec

cooling curvt second derivative

Time (sec)

Figure 5.10a : The cooling curve and associated derivative curve for a 3 19.2 alloy having 0.36% wt iron.

Figure 5.lOb : The microstructure of the 3 19.2 alloy according to the cooling c w e above. 1 .) P-phase, 2.) silicon, 3 .) & Cu-Al-Si temary eutectic & 4.) a-aluminurn

CHAPTER FWE : RESULTS OF THERMAL ANALYSIS 47

Cooling Curve & Associated Second Derivative for 3 19.2 alloy (trial two) : Reference cooling rate = O. 1 O°C/sec

-

-

450 ' 1 I I I 1 I I -0.02 O 200 400 600 800 1,000 1,200 1,400 1,600

Time (sec)

Figure 5.1 la : The cooling cuve and associated derivative c w e for a 3 19.2 alloy having approximately 0.92% &t iron.

Figure 5.1 1 b : The microstructure of the 3 19.2 alloy according to the cooling cuve above. 1 .) $-phase, 2.) silicon, 3 .) Al 2Cu-Al-Si temary eutectic & 4.) a-aluminum

CHAPTER FTVE : RESULTS OF THERMAL ANALYSIS 48

Cooling Curve & Associated Second Derivative for 319.2 alloy (trial two) : Reference cooling rate = O. 10°C/sec

Time (sec)

Figure 5.12a : The cooling curve and associated derivative c w e for 3 19.2 alloy having 1.25% wt iron.

Figure 5.12b : The microstructure of the 3 19.2 alloy according to the cooling cuve above. 1 .) B-phase, 2.) silicon, 3 .) & Cu-Al-Si temary eutectic & 4.) aaluminum


Cooling Curve & Associated second Derivative for A4 13 Alloy

Figure 0.93%

400 time (sec)

5.13a: The cooling m e and associated derivative curve for a A4 13 al wt iron.

loy having

Figure 5 curve ab

.13b : The microstructure of the A413 alloy solidified according to the ove . 1.) a-phase, 2.) silicon & 3.) a-duminum.

cooling


Cooling Curve & Associated Second Derivative for A4 13 Alloy

450 ' I 1 I

-0.02 O 200 400 600 800

time (sec)

Figure 5.14a : The cooling curve and associated derivative cume for a A413 alloy having 1.4% wt iron,

Figure 5.14b : The microstructure of the A413 alloy solidified according to the cooling cuve above . 1.) P-phase, 2.) silicon & 3 .) a-aluminum.


Cooling Curve & Associated Derivative for A4 13 Alloy

400 time (sec)

Figure S. 15a : The cooling curve and Associated denvative Curve for a A413 alloy having 2.01% wt iron.

Figure 5.15b : The microstructure of the A4 13 alloy solidified according to the cooling curve above . 1.) P-phase, 2.) silicon & 3.) a-aluminum.

CHAPTER FIW : RESULTS OF THERMAL ANALYSTS 52

a mainiy alurninum-silicon eutectic structure with both AI$eSi and A.i,,(Mn,Fe),S&

intermetallics. The morphology of the Al15(Mn,Fe)& phase in figure S. 13b is chinese

script. Figure 5.14a shows a A413 alloy coohg curve with an iron content adjusted to

1.4% wt. Two new thermal anomalies appear on the cooling curve. The fint one is

believed due to the formation of Al&Si. The second anomaly is associated with the

formation of prirnary a-alurninum. This thermal effect is believed to be due to prirnary

a-aluminurn since the other alternative is primary Silicon, but no polyhedral siLicon

crystais are observeci in the microstnicture seen in figure 5.14b. Figure 5.15a shows the

cooiing curve for the A413 alloy having 2.0% wt iron. The duration of the *F&i

anomaly has increased as a result of the increase in iron content along with an increase

in the formation time of the primary a-aluminum. In the associated micrograph in figure

5.15b, the Ai,,(Mn,Fe)3Si, phase is present along with very long AiPeSi needles. The

morphology of the Ai,,(Mn,Fe)3Si2 is different from the previous two A413 samples in

that it is more star-like.

CHAPTER FiVE : RESULTS OF THERMAL ANALYSE5 53

5.2 Quantification of Fe-bearing Intermetallies Within this section an explanation is given of how the second denvative curve, produced

from the coohg curve, is to be interpreted and how it can be used to quanti@ formation

temperatures and duration of resolvable A15FeSi and alurninum-siliwn eutectic thema1

anodes.

5.2.1 A&FeSi Thermal Signature

The formation of A1,FeSi results in the release of latent heat of fusion in a solidimg

alloy leading to an anomaly (Le. irregularity on the cooling curve). The rwlution of

such a thermal anomaly is dependent on whether the phase foms during a time not

coincidentai with the major phase reactions (Le. pnmary a-aiuminum or eutectic

growth). For lower Fe contents the formation of A15FeSi does occur, but its appearana

on the cooling curve is usually obscured by the eutectic plateau. At higher Fe contents,

according to the phase diagram of figure 2. lb, the crystallization tempemure of AsFeSi

occurs at temperatures higher than that of the AI-Si eutectic. As a consequence, the

beginning of the thermal signature of this phase fonnation becomes resolvable. As iron

content gradually increases, the start temperature of this phase will continue to increase

until it coincides with the thermal signature of the primary arrest, at which point it will

again be unresolvable. The total time duraton of this thermal signature (time parameter

as seen in figure 5.16) does not reflect the total A&FeSi that forms, since the formation

of this phase continues until the end of alloy solidification. However the duration of the

thermal anornaly does correlate with the amount of Fe in the melt, and thus it is

reasonable to quantify the Fe content by measuring some parameter associated with this

anomalous effect.

5.2.2 Quantifying Fe via Thermal Analysis

No r d attempts have been made to date quanti@ to Fe contents via thermal analysis.

Usuaily some off-line analytical technique such as spectmchemical analysis is used. For

CHAPTER FWE : RESULTS OF THERMAL ANALYSE 54

an on-line method thermal analysis is a more practicai approach. Two methodologies to

relate A15FeSi amounts to the cooling cume are given.

5.2.2.1 Temperature Method

Ananthanarayanan et al'833 have suggested that the crystailization temperature of Al$Wi

could be determined from the peak signature of the fint derivative curve, then wrrelated

with the point on the cooling curve and extrapolated back to the vertical axis

(temperare aris). It was found that the cooling rate and degree of superheat affected

the formation temperature of the AlsFeSi thermal signature as determined from this

method, and no attempt was made to use it for varying Fe contents. While this method

could be useful, it relies on themocouple accuracy.

An alternative to indicating the temperatures of heat effects on cooling curves wiii

be developed in this thesis. This will involve the use of maxima on the second derivative

curve. The fint maximum of the dq/dt? curve associated with the A1,FeSi thermal

anomaly cm be used to determine the resolvable fornimion temperature, labelied TH,,

on figure 5.16. This is determined by correlation from the d?/dt2 maximum, to the

cooling curve and then extrapolated to the vertical axis. It will be assumed that this

temperature is very close to the equilibnum nucleation temperature of AisFeSi

intermetallic at very slow cwling rates. The end temperature for the A1,FeSi heat effect

is determined from the next maximum encountered on the d'T/d? curve, labelied as Tb,.

in figure 5.16. This maximum is actuaily a response to the change in cooling rate

associated with the beginning of the eutectic plateau. The actual end temperature for

A1,FeSi formation is not resolvable since the heat effect continues into eutectic

solidification. Thus, this temperature is really a pseudo-end remperafure for the A1,FeSi

intermetallic.

Another temperature parameter of interest is TM., the start temperature of primary

aluminum solidification in high iron 413 alloy. Studying the variation of T, is important

since it wiil help explain what may be happening to the aiioy when iron content

increases. This is illustrated on figure 5.17.

CHAPTER FIVE : R E S W OF THERMAL ANALYSIS 55

5.2.2.2 T m e Method

A rarely considered aspect of thermal analysis is to measure the time it takes for a phase

to grow. The text by Tenekedjiev et alz made an atkmpt to measure a time parameter

for the crystallization of the eutectic with modifier and no modifier. The method used to

determine the start and end times for phase formation was to locate the inflections on the

cooling curve before and after the thermal anomaly for the phase in question. These

inflection points were located by drawing tangentid lines, or slopes, before and after the

suspected inflection point. The intersection of the tangentid lines was then extrapolated

vertically to mark the inflection or the start and end times for the phase reaction. The

major problems associated with this methodology are that determination of inflection

points by such a manual method can lead to values subject to personal interpretation, and

some inflections are sirnply very difficult to detect in this way.

An alternative to determining start and end times is to find local maxima of the

d?Udt2 curve associated with the inflections of the phase in question. This concept has

been used to identify the beginning of eutectic growth in cast iron thermal analysis", and

is ihstrated in figure 5.16 for an A1,FeSi thermal signature in 356 alloy. The sran time

is deterrnined by the position of the first maximum on the d2~ /d? curve near the vicinity

of the initial A1,FeSi thermal anomaly. The actual end time cannot be determined since

solidification of A1,FeSi continues into eutectic solidification. There is a pronounced

maximum on the d2~/dt2 curve at the beginning of eutectic solidification, and this is used

as an arbitrary end point for the AISFeSi thermal anomaly. This end time is then really

a pseudo-end rime, and the whole observable AisFeSi thermal anomaly duration is in

effect an apparent rime paramefer. For the eutectic alloy (413) studied, the A1,FeSi that

is resolvable on cooling curves grows in pure liquid as a primary phase. This is shown

in figure 5.17.

5.2.2.3 Tome & Temperature Method for Eutectic Soliciifkation

It wiU be shown that the duration of the aluminum-silicon eutectic plateau on cooling

curves is affected by the presence of substantial iron contents if A1,FeSi crystallizes as

a primary phase. The start and end times are detennined by the maxima on the second

CHAPER FIVE : RESULTS OF THERMAL ANALYSB 56

derivative curve. The starr time of the eutectic solidification is show in figure 5.18 and

is the same as the AisFeSi pseudoend time. The eutectic start temperature is also

determined by this infiection and cm be extrapolated to the vertical axis h m the cooling

cume. This is labeiieû Tm,-. The is also the pseudoend time or temperature fi,) for the A1,FeSi as describeci in the previous sections. The end time for the eutectic

solidification occurs at the local maximum in the second derivative curve that

corresponds to the inflection associated with the beginning of the post-eutectic phase

reaction. This reaction is the Ai,Cu-Ai-Si eutectic for 319 aUoys, or the Mg,Si-Al-Si

eutectic for the 356 alloy. Extrapolation from the cooling curve back to the vertical axis

will indicate the temperature where eutectic freezing stops, labelled T, ,. There exists

no such posteutectic reaction for 413 alloy, thus some alternative method to determine

the eutectic end time and temperature is needed. In cast iron thermal analysis the end of

eutectic freezing results in a change from a zero cooling rate to a finite cooling rate. This

results in a negative peak on the second denvative curve4'. This absolute minimum could

be used as an indicator of the end of eutectic freezing for the 413 alloy, as illustrated in

figures 5.19 and 5.20.

When the iron content of the 413 alloy increases, the cooling curve changes from

one with a single eutectic plateau to a more complex curve having three distinct heat

effects: Al,FeSi, a-aluminum and the Al& eutectic plateau. Figure 5.20 shows the

cooling curve for a high iron 413 ailoy. The eutectic stan time is dekrmined to be the

end of primary a-aluminum growth (eg. local maximum of the dq/df) instead of the

juncture between pure Iiquid cooling and eutectic start temperature as Ken for low iron

413 in figure 5.19.

Due to the slow cooling rate used in this research, the minima in the second

derivative curve which indicate the end of eutectic freezing are often quite broad making

it difficult to find a distinct minimum in the d?Udf curve. In this case, a local maximum

in the fourth derivative curve can be used to signal the end of eutectic fieezing as shown

in figure 5.2 la. Figure 5.2 1b shows that when the fourth derivative is at a maximum the

second denvative is at a minimum. The advantage of using the fourth denvative is that

it is easier to observe than a minimum in the second denvative.

Detennination of Apparent Time Parameter Iron-b&ng Intermetallic of 356 & 3 19 Alloys

300 400 time (sec)

Figue 5.16 : The maxima of the second denvative are indicated in this figure. The first maximum occurs before the b-phase thermal anomaiy (1) and the second before the beginning of eutectic growth (2). These maxima represent the start and end time and temperature based parameten of the P-phase anomaly.

CHAPTER F~VE : RESULTS OF THERMAL ANALYSIS sa

Deternination of Apparent Time Parameter Iron-bearing Intermetallic & Primary Aluminurn of 4 13 Ailoy

100 150 time (sec)

Figure 5.17 : The maxima of the second derivative are indicated on the figure. The three major peaks observed correspond to the beginning of P-phase growth (1 .), beginning of a-aiuminum primary growth (2.) and the beginning of Al-Si eutectic growth (3.). The temperature based parameten measured are : Tpr-, for the begiming of B-phase growth, TM., for the beginning of primary aiuminum and Taam. for the beginning of eutectic growth.

R FIVE : RESULTS OF THERMAL ANALYSE 59

Determination of Eutectic Solidification Time Eutectic Solidification Tirne for 3 19 & 356 AUoy

Second derivative I l - I I 1 1 I I 1 I I ,

300 400 500 600 700 800 906 1

time (sec)

Figure 5.18 : The maxima of the second denvative can be used to indicate the begi nning of eutectic growth (1 .) and the beginning of the post-eutectic reaction (2 .). The time duration between these two peaks cm be used to mesure eutectic time duration. Temperature parameters meanired are Taa nr for the beginning of Ai-Si eutectic growth and Tclncnd for the beginning of post-euteaic growth, also considered the end temperature of the main Ai-Si eutectic formation.

CHAPTER FNE : RESULTS OF THERMAL ANALYSE

Eutectic tirne detemination for low Iron 413 for A41 3 Alloy : Fe content 4 . 9 % ~

time (sec)

Figure 5.19 : For low iron 413 ailoys the maximum of the second derivative is labelled as (1 .). This maximum reflects the change in cooling rate fi-om pure liquid cool ing to eutectic cooling. The minimum in the second derivative corresponds to the end of eutectic fkeezing (2.).

CHAPTER FïVE : RESULTS OF THERMAL ANALYSIS 61

Eutectic time detennination for high Iron 413 for A41 3 Alloy : Fe content = 2 . 2 % ~

time (sec)

Primary aluminum

P-phase growth

Figure 5.20: For high iron 413 alloy three maxima of the second derivative cuwe occur due to the presence of P-phase and primary a-aluminum heat effects. For these ailoys the start time and temperahire for eutectic fieezing begins after the primary aluminum heat effect, not aftu pure Iiquid cooling. The end temperature and time for the main Ai-Si eutectic is the minimum in the second derivative.

Cooling Curve with Fourth Derivative A413 Alloy

I I 1 1 I J 200 400 600 800

time (sec) Figure 5.2 la: Cooling curve of a 413 alloy with fourth denvative. Notice the maximum near the end of eutectic plateau.

Cornparison of Second & Fourth Derivative A413 Alloy

a50 500 550 600 650 time (sec)

Figure 5.2 1 b : This figure repnsents the dashed box in figure 5.2 la. The minimum of the second denvative corresponds to the local maximum of the fourth derivative.

C-R FIVE : RESULTS OF THERMAL ANALYSTS 63

5.3 Tabular Results of T h e d Analysis Data

The start and end temperatures, and duration of the important thermal anomalies

measured h m the cooiing curves by the rnethods just describecl are presented in the

foliowing tables.

Table 5.1 : Temperature and Time Based Parameten for A1,FeSi intermetallic in A356 AUoy. Reference Cooling Rate = 0.100Clsec

Target

ÇHAPTER FlVE : RESULTS OF THERMAL ANALYSE

Table 5.2 : Temperature and Time Based Parameters for AipeSi Intermetallic in A356 Alloy. Reference Cwiing Rate = 0.4X/sec

CHAPTER E;IVE : RESULTS OF THERMAL ANALYSIS 65

Table 5.3 : Temperature and T i e Based Parameters for AlScSi Intermetallic in 3 19.2 Ailoy. Reference cooling rate = O. 1O"CIsec

CHAPTER FlVE : RESULTS OF THERMAL ANALYSE

Table 5.4 : Temperature and Time Based Parameten for the AlSeSi ïntermaallic in 3 19.2 Alloy . Reference cooling rate = O. 1O"CIsec

Trial Two

CHAPTER FLVE : RESULTS OF THERMAL ANALYSE 67

Table 5.5 : Eutectic time parameters detezmined fkom cooling curves of hypoeutectic (356 and 319) alloys studied. The start and end temperatures for these alloys were found to be invariant m increasing iron content, and are not presented here. The target iron levels are given in tables 5.1, 5.3 and 5.4.

3 19.2 Alloy (O. 1VCIsec) trial one

3 19.2 Alloy (O. 1CPClsec) trial two

% wt. Fe Time (sec)

CHAPTER FNE : RESULTS OF THERMAL ANALYSE 68

Table 5.6 : Time Based Parameters of the AIPeSi intermetailic, or-alurninum and the Ai-Si eutectic for A4 13. The insulated cup method was used.

lWFi : RESULTS OF THERMAL ANALYSE 69

Table 5.7 : Temperature Based Parameters of the AisFeSi intmetallic, a-aluminum and AI-Si eutectic for A413 Alloy. The insulated cup method was used.

* Eutectic end time deterrnined via local maximum of the dWdP

CHAPTER FIVE : RESULTS OF THERMAL MALYSIS 7Q

5.3.1 General Trends Observed in Tabled ResuII

According to the data in tables 5.1, 5.2, 5.3, 5.4 and 5.6 the formation temperature

and time duration for the Al,FeSi thermal signature increase with increasing iron

content for the three aüoys studied. An interesting obsewation is the cunsistency of the

rneasured parameters from repeats performed. It is clear from tables 5.1, 5.2, 5.3 and

5.4 that any slight variation of these parameters for a given iron content seems to be

related to the thermal analysis sample mass. For example, in table 5.3, the three thermal

anaiysis sarnples having an iron content of 0.74 96 wt had sarnple masses of 227.18,

228.1g and 244.1g. The sample having the mass of 244.1g had a longer A1,FeSi time

duration, and according to table 5.5, a longer eutectic time duration, than did the other

two samples. At the same time, the resolvable formation temperature for the AlfieSi

thermal signature and eutectic plateau was repeatable within the standard accepted

thermocouple error of k2OC. This is an indication that temperature based parameters do

not show the same dependency on sample masr. The observation of time parameter

repeatability being sample mass dependent will explain the slight scatter of repeats in

subsequent figures showing both A1,FeSi and Al-Si eutectic duration versus iron content.

5.4 Image Analysis Ail thermal analysis sarnples were cut near the tip of the thermocouple and polished to

be examined via image analysis. The purpose of image analysis is to establish a

relationship between the percent area that A1,FeSi occupies within the microstmcture and

the parameters of the A1,FeSi thermal anomaly as determined from the thermal analysis

curves. The error bars seen in the following figures which present image analysis results

represent the standard deviation of the mean percentage area from 10 to 15 frames taken

for each metdlographic sarnple. The best fitting linear regression lines, using FreelanceTM

for Windows software, were used in these figures to help illustrate the correlation of

apparent time parameter versus percent area intermetallics in the microstmcture.

CHAPTER FLVE : RESULTS OF THERMAL ANALYSIS 71

5.4.1 Image Analysis of A356 Alloy

Figures 5.22 and 5.23 show plots of percent area A1peSi phase in the microstructure of

the thermal analysis sarnples versus apparent time parameter for A356 alloy using the

insulated cup method and the sand method, respectively. It *ui be obsewed that the best

fitting line for the insulated cup method is steeper than for the sand method. The reason

for this may be due to the fact that at the slower cooling rate the AhFeSi phase had more

time to grow before the liquid was nch enough in silicon to achieve eutectic growth. For

example, in samples having 1.25 % wt Fe (tables 5.1 & 5.2) the 4FeSi phase grew for

600 seconds in the insulated cup and 100 seconds in the sand mold.

5.4.2 Image Analysis for 319.2 AUoy

Figures 5.24 and 5.25 show the percent area AI,FeSi in the microstructure of the thermal

analysis sarnples versus the apparent time parameters for the two repeated expenments,

trials one and two. For tnal one the average sample size was 233 * 10 grams and for trial

two 208*5 grams. The slopes of the best fitting iine using the mean values for both

figures are about the same. The constancy of the mean values for triai two is slightly

better than for trial one for a given iron content, due to the fact that sample mass was

more consistent.

5.4.3 Image Analysis for A413 AUoy

Figure 5.26 shows the percent area that Aimi occupies in the microstructure versus

the apparent time of the A15FeSi themal anomaly. The A1,FeSi thermal anomaly was

only observable at high iron levels. As a consequence, results from only the last seven

cooling curves obtained are shown in this figure. Figure 5.27 shows the percent area of

the A15FeSi intermetallic versus the time for eutectic solidification. The eutectic duration

is reduced as the arnount of the A1,FeSi phases increases within the microstmcture. It is

possible that the reduction of the eutectic duration occurs as siiicon is depleted from the

rnelt in the formation of pre-eutectic A&FeSi.

RESW.TS OF THERMAL ANALYSE 72

Figure 5.22: % Area of B-Phase in Microstructure vs. Apparent Time A3 56 Aiioy : Referieirce coofmg rate : O. 1 O'Usec

1, --

.-

1 I I I I O 200 400 600 800 1,000

Apparent tirne parameter of &phase (sec)

Figure 5.23 : ?/o Area B-Phase in Microstructure vs. Apparent Time A356 Alloy : Reference cooling rate = 0.45'C/sec

60 80 100 Apparent time parameter of &phase (sec)

Figure 5.24 : % Area of &Phase in Microstructure vs. Apparent Tirne

1 O0 200 300 400 500 Apparent T i Parameter of beta-Phase (sec)

Figure 5.25 : % Area of B-Phase in Microstructure vs. Apparent Time

O 1 O0 200 300 400 500 Apparent Timc Pararneter of bcta-phase (sec)

CIfAPTER FïVE : RESULTS OF THERMAL ANALYSE 74

Figure 5.26 :% Area 13-Phase in Microstructure vs. Apparent Time A413 AUoy : Insulated cup method

2 ,

O - . 1 1 I l t . 50 60 70 80 90 100 110

Apparent tirne parameter of &phase (sec)

Figure 5.27 : % Area &Phase in Microstructure vs. Eutectic Duration A413 AUoy : Insulated cup method

O 12

r d

3 ,O

t -- '

i!

f * - o. Q

8 4 : s

2

O

--

-- -- --

1 I 1

700 750 800 850 900 Eutcctic time parameter (sec)

U?.TS OF THERMAL ANALYSIS 75

5.5 Cornparison of Formation Temperature & Apparent Tirne

Parameter for Quantifying Iron content of

AIuminumSiiicon Ailoys

The overall objective of this project is to estabiish a simple and reliable methodology to

quanti@ the iron content of the melt by meashg a particular parameter of the W e S i

thermal anomaly. As indicated in the previous sections, both the apparent time and the

resolvable formation temperature of the A15FeSi were measured h m all the cooling

c w e s wmpleted. Thus it is possible to relate the Al,FeSi formation temperature and

time duration to the alloy iron content.

5.5.1 A356 W o y

Figures 5.28 and 5.29 show plots of the formation temperature of A.i$eSi versus the

weight percent iron for the insulated cup methoâ and sand mold method respectively. The

slopes of the b a t fitting line for the two different coolhg rates are 40 and 36.6 " C M wt

Fe respectively . The lower limit of detectability is 0.54 % wt Fe for the insulated mold

method, but shifted up to 0.74% wt Fe for the sand mold method. Both are indicated on

the figures. While there is no real difference in dope, the curve is shifted to the right

(i.e. to higher iron content) for the data taken h m the sand mold method. To use this

method of measuring formation temperature to gauge iron content requires a strict

adherence to a particular thermal analysis setup having a repeatable reference cooling

rate.

Figures 5.30 and 5.3 1 show the plot of A15FeSi time duration versus weight

percent iron for the insulated cup method and sand mold method respectively. The slopes

of the best fitting lines are different for the experiments using the two different cooling

rates. The slope for the slower cooling rate is approximately 950 seconds/wt% Fe, while

the faster cooiing rate yields 375 seconds/wt% Fe. This differenœ in slope is due to the

fact that at faster cooling rate the Al,FeSi has less time to form at a given iron content.

For ewnple, an alloy having 1.25 wt% iron, will allow approximately 100 seconds for


Figure 5.28 : Formation Temperature of B-Phase vs. Iron Content A356 Alloy : Rcfmœ cooling rate = 0.1O'Uscc

1 6 Lower b i t of iron detection

0.8 1 1.2 1.4 Percent weight iron

Figure 5.29 : Formation Temperature of B-Phase vs. Iron Content A356 Alloy : Reference coohg rate = OAS'Usec

Lower M t of iron detection

0.8 1 1.2 Percent weight iron

RESUI.TS OF THERMAL ANALYSIS 77

Figure 5.30 : Apparent Time of &Phase Growth vs. Iron Content A356 Aiioy : Referenece cooling rate = 0.1 OaC/sec

V

0.4 0.6 0.8 1 1.2 1.4 Percent weight iron

Figure 5.3 1 : Apparent T h e Parameter for &Phase Growth vs. Iron Content A356 Alloy : Refcraiece cooling rate = 0.45'Usec

Percent weight iron

CHAPTER FiVE : RESULTS OF THERMAL ANALYSIS 78

Figure 5.32 : Formation Temperature of B-Phase vs. Iron Content Trial One : 3 19.2 AUoy : Refcfencc cooling ratt = 0.1 O' Usec

Figure 5.33 : Formation Temperature of B-Phase vs. Iron Content Triai Two : 3 192 Alloy : Reference coohg rate = O. 1 O'Usec

0.6 0.8 1 1.2 1.4 Percent weight iron

605

600 O - 595

9 59, E ,585

580 G

.O 575

570 565

560

:

1 Lower Mt of iron detection

/q v

b 1 I I I

0.6 0.8 1 1.2 1.4 Percent weight iron

W.TS OF THERMAL ANALYSIS 79

Figure 5.34 : Apparent Time for B-Phase Growth vs. Iron Content 319.2 Alloy : Trial ûnc

rc 300

b E

L zoo

0.8 1 1.2 Percent weight iron

Figure 5.35 : Apparent Time of B-Phase Growth vs. Iron Content 3 19.2 AUoy : Trial Two


Figure 5.36 : Formation Tem~erature of &Phase vs. Iron Content A

A41 3 Ailoy : lnsulated cup method

2 2.1 2.2 2.3 2.4 2.2 Percent weight iron

Figure 5.37 : Formation Temperature of Primary Aluminurn vs. Iron Content A4 13 Alloy : Insulated cup method

1 -2 1.4 1.6 1.8 2 2.2 2.4 2.6 Percent weight iron

CHAFïER FIVE : RESULTS OF THERMAL ANALYSE 81

A&FeSi growth at the faster cooiing rate, but 650 seconds at the slower cooling rate. It

becornes clear that time based parameters are strongly ïnfluenced by cooling rate.

5.5.2 319.2 AUoy

Figures 5.32 and 5.33 show the plot of formation temperature versus iron content. The

dopes of the best fitting lines are 50 and 53 OC/%wt Fe for trials one and two

respectively. The lower Limits of iron detection using this method are indicated on the

figures. Figures 5.34 and 5.35 show the plot of apparent time parameter for A&FeSi

growth versus iron content. The slopes of the best fitting lines are 516 seconds/wt% Fe

and 526 seconds/wt% Fe for trials one and two respectively. Both W s demonstxate

repeatabiiity since the best fitting lines are reasonably close, giving confidence that such

an approach is a reliable method for predicting iron content.

5.5.3 A413 Aiioy

Figure 5.36 shows the plot of Al,FeSi formation temperature versus iron content. The

A15FeSi thermal anomaly was not resolvable at iron contents less than 1.8 96. At 1.8 %

wt iron the AisFeSi anomaly finaily appeared. As a consequence, this method of

quantifjhg iron is not very successful in this eutectic aUoy . Figure 5.37 shows the sarne

variation of iron content, but plotted against the formation temperature of the primary

a-aluminum thermal anomaly. This anomaly appeared at 1.1 % wt iron and then gradually

increased in formation temperature with increasing iron content.

CfFAPTER FIVE : RESULTS OF THERMAL ANALYSE 82

5.6 Effect of Iron on Eutectic Time hiration

The eutectic duration was observed to decrease with increasing iron content as shown in

table 5.5 for the two hypoeutectic alloys studied, and in table 5.7 for the eutectic ailoy

studied.

5.6.1 Aiioy A356

The decrease in the duration of eutectic solidification was found for the thermal analysis

test samples using the insulated cup method (O. lVC/sec). The rand mold method did not

result in a noticeable decrease in eutectic duration with increasing iron content, indicating

that this is an effect that becornes obsenable only at very slow cooling rates. Figure 5 -38

shows the eutectic solidification time for A356 alloy versus iron content. The best fitting

line through the data is dashed and not solid because of the scatter in the data. It is

believed that the scatter of the repeats of thermal analysis samples for a given iron

content reflects the fact that sample size was not always constant; however, sample mass

was not measured for this experiment.

5.6.2 AUoy A413

Figure 5.39 shows the plot of eutectic solidification time of 413 alloy versus iron content.

For the hil iron range studied in this experiment the eutectic time appears to have

decreased as iron content increases. As in the case with the A356 alloy the scatter of the

&ta most likely is due to variation of the sample mass itself. The mass of the thermal

analysis samples were measured before cutting for metallographic preparation, and these

values are shown on the figure. As can be seen, a relationship between the mass of the

sample and eutectic time for a given iron content exists. Longer times are found in

samples of larger mass with relatively srnail changes in mass resulting in very

significantly longer eutectic ti me durations.


5.6.3 AUoy 319.2

The thermal analysis experiment for this alloy was repeated twice to estabiish the

inte- of the data obtained for increasing iron contents. For trial two, particular

attention to controliïng sample mass was taken. It was found during aial one experimen~

that cooling curves were on the whole of differing durations. For example, samples

which were noticeauly larger in size resulted in a longer eutectic solidification time. The

driving force to perform triai two was to establish if controlling and recording the actual

sarnple mass had any effect on the duration of the eutectic freezing. If this was so, then

it could explain the scatter in results seen in pewious thermal analysis experiments.

Cornparison between trails one and two, seen in figures 5.40 and 5.41

respectively, shows the effect of sample size consistency. Plotting eutectic time versus

iron content yeilds almost no trend for trial one. The measured mass of the thennal

analysis sarnples before any cutting are given in the figure. The overall mean samp1e

mass was 233f 10 grams. As seen in figure 5.41, better sample mass control yeilds

consistency between repeats, and a more noticeable trend is found. The overail mean

sample size in this case was 208 &5 grams. The best fitting line is marked by a dashed

line. The decreasing eutectic time with increasing iron content is consistent with what

was observed for the A356 and A413 alloys. Clearly the variation in the eutectic time

parameter indicates that using a time based parameter requires strict adherence to sample

size for a given thermal analysis setup.

5.6.4 Eutectic Formation Temperatures

The s ta r t and end temperatures of the eutectic reaction in 3 19.2 (trial two) alloy were

detennined by using the local maxima of the dqld? curve before the eutectic

solidification and before the post-eutectic reaction (end-time). In figure 5.42 these are

plotted versus iron content. The start and end temperatures for the eritectic of the A413

alloy were determined by the locai maximum of the dq/dt2 curve before the eutectic

solidification, and the minimum of the dq/dt2 cuve at the end of the eutectic. These are

plotted in figure 5.43 as a function of iron content. It would appear that these

tempemtures do not change with increasing iron wntent despite the fact that eutectic

Figure 5.38 : Eutectic Duration vs. Iron Content A356 Alloy : Reference coohg rate = O. 1 OeC/sec

Percent weight iron

Figure 5.39 : Eutectic Duration vs. Iron Content A41 3 AUoy : Insulated cup method

1 Each data point includes the sample mass 600 ' I I 1 t

1 1.5 2 2.5 Percent weight iron

C-R FIVE : RESULTS OF THERMAL ANALYSIS 85

Figure 5.40 : Eutectic Duration vs. Iron Content 3 1 9.2 Alloy : Trial One : Reference cooling rate = O. 1 O' C/sec

* 1 Each daia point includes the -pie m a s in gr- 550 ' I 1 1 l I I

0.2 0.4 0.6 0.8 1 1-2 1.4 Percent weight iron

Figure 5.4 1 : Eutectic Duration vs. Iron Content 3 19.2 Alloy : Trial Two : Reference cooling rate = O. 10°C/sec

- - - 0.2 0.4 0.6 0-8 I 1.2 1.4

Percent weight bon

U .C. - 0

3 600 d

-

Each data point includes sample mass in gram I I I lq8gm 550 I


Figure 5.42 : Eutectic Start & End Temperatures 3 1 9.2 Alloy (trail two)

0.6 0.8 1 Percent weight iron

Figure 5.43 : Eutectic Start & End Temperatures A4 1 3 Alloy

Percent weight iron

CHAFïER FIVE : RESULTS OF THERMAL ANALYSE 87

Figure 5.44 :Determination of the End of Eutectic Solidification

1 1.5 2 h n content (Yo wt)

Figure 5.45 : Detemination of the End of Eutectic Solidification Iacnl Msxirnmi of lbt F o d l JkhmivC ., 564

'[ 562

5 560 CE O aSs8 C; 556 û - 554 f g 552 e + 550

- - 8 .

- - - - - D

8 . .

y - m - -

I 1 4

O 5 1 15 2 2.5 I n n i content (% wt)

duration does decrase over the same iron composition range.

It was pointed out earlier that the fourth derivative wuld be of asoistance in

detaminiag the actuai end time of the eutectic soiidification in the 413 ailoy. As shown

In figure 5.21a and b the maximum of the fourth derivative has the advantage of king

easier to see. Figures 5.44 and 5.45 present a cornparison of the end temperature of

eutectic freezing for 413 alloy detennined from the second and fourth derivatives

respectively. The scatter in both figures is due to thennal noise, caused by convection

in the melts. Within the limits of the present expriment, these two methods yield

essentiîlly the same results for the end of eutectic solidification for 413 aiioy.

SIX : DISCUSSION OF RESULTS 89

Chapter Six

Discussion of Results

6.0 Discussion of Results The coolhg curves aquired for this thesis wili now be correlateci to the hown portion

of the Al-Fe-Si temary phase diagram. From this, a discussion on how feasible it is to

use thermai analysis to quanti@ Fe content wül be made.

6.0.1 Hypoeutectic AlSi Alloys

For the two hypoeutectic aluminum-Silicon aUoys studied (A356 & 319.2) the A&FeSi

t h e d signatures that where detectable on cooling curves obtained for this thesis

indicate that these particular intermetallics grow within the interdendritic regions dunng

solidification when iron contents are approximately 0.6% wt or higher. The resolvable

formation temperature of this intennetallic should ideally be close to the equilibrium

formation temperatures when using the insulating cup method (O. lO"C/sec). Thus

cornparison to the appropriate liquidus sUTface of the Al-Fe-Si phase diagram wouid be

redistic. For the A356 (7.2% wt siricon) alloy this can be done using the aluminum-iron

pseudo-binary havhg a constant of 8.0% wt Silicon. The slope of the liquidus iine for

the Al+B+liquid field in figure 6.0 is 38.3*C/%wt Fe while the dope of the best fitting

line of the fonnation temperature of the AlPe!Si with increasing iron content is 40"CI%wt

Fe as seen in figure 5.27. With the formation temperature data p l a d on figure 6.0, the

agreement with the liquidus line appears to be very good.

For the A356 aiioy, using the sand mold method (O.4S0Clsec), the slope of the

formation temperature line for this intermetallic is 36.6OC196 wt Fe. This iine is plotted

on the same pseudo-binary in figure 6.1. While the slope is in good agreement with that

CHAPTER SIX : DISCUSSION OF RESULTS

Figure 6.0 :Formation Temperature of B-Phase vs. Iron Content A356 Aiioy : Reference cooiing rate = 0.1 0 C/sec

Percent iron content

Figure 6.1 Formation Temperature of B-Phase vs. Iron Content A356 Alloy : Reference cooling rate = 0.45 C/sec

Percent weight iron

605 - /

Û 600 595

E g 580 .' E 575

- Al + Liquid

I Ai + B + Liquid

- c-" 570 -

CHAFTER SIX : DISCUSSION OF RESUtTS 91

Figure 6.2 : Formation Temperature of B-Phase vs. lion Content Trial One : 3 19.2 AUoy : Reference cooling rate = 0.10 U s c

620 L

Percent weight iron

Figure 6.3 : Formation Temperature of B-Phase vs. Iron Content Triai ?tvo : 3 19.2 Ailoy : Reference coolhg rate = 0.10 Usec

620

Percent weight iron

CHAPTER SIX : DISCUSSTON OF RESULTS 92

of the liquidus, the achial data points are shifted to the right. This is a result of the faster

cooling rate (an increase of 4.5 times) which does not allow for equilibrium formation

of AlSFeSi. As a consequeme, the lower limit of deteetab- using the insulated cup

method was increased from 0.54% wt Fe to 0.75% wt Fe for the sand mold method.

However, at the same time the upper limit of detectability increased from 1.2 % wt to

1.4% wt Fe.

For the 319.2 alloy the slope of the Al5FeSi resolvable formation temperature

versus iron content is S O T / % wt Fe and 53"C/% wt Fe for trials one and two

respectively. Their slopes agree closely with the slope of the liquidus of the

Ai+p +liquid field reproduced for a constant vertical Silicon content of 6%wt, shown in

figures 6.2 and 6.3. However, it is clear that the actual data is shifted to the right and

downward with respect to the actual liquidus. The copper content which is of the order

of 4 wt% could explain the shiftllig that has occurred in both figures. This shifting to the

right and downward couid explain why the lower limit of iron detectability was higher

for 319 alloy than for 356 alloy, despite the fact that 319 alloy has a lower eutectic

temperature which will not obscure the A.&FeSi thermal signature.

The repeatability of the two resolvable formation temperature profiles for 3 19

aUoy (trails one and two) detennined by the variability of the Ai$& thermal anomaly

is quite good. This indicates that imn content measured by the thermal analysis method

could be reliable.

The duration of AisFeSi formation also increased with increasing iron content

within the detectability range, for the A356 and 319.2 alloys. The increase in time

duration with increasing iron could be an alternative to measuring the formation

temperature, and would have the advantage of being thermocouple independent. However

the time duration method has more scatter for repeats of a given iron content, probably

due to süght inconsistencies in sample mas . Examples of this can be seen in tables 5.3,

5.4 and 5.6. Improved thermal analysis equipment which can control sample mass and

heat extraction, would make time parameters far more reliable then measuring

temperature.

It was observed in both hypoeutectic aiioys that, while start and end temperatures

CHAPTER SIX : DISCUSSION OF RESULTS 93

for the eutectic did not change with increasing iron, the duration of the eutectic did. This

result is most W y due to the fact that siiicon is depleted before eutectic freezing in the

formation of prirnary Al,FeSi, resulting in less silicon available for the actuai aluminum-

silicon eutectic. The change in eutectic time with increasing iron is less dramatic than is

the A15FeSi heat e&t, and thus it could never realisticaily assist in iron determination

in the melt.

6.0.1 Eutectic AI-Si AUoy

In the A413 aüoy the behaviour of the coolhg curves with increasing iron content

is very cornplex. Figures 5.12a-5.15a show the cooling curves of A413 dloy with

increasing iron content. It was found that the duration of the eutectic plateau, the AisFeSi

heat effect, and the primary a-alurninum heat effect changed dramaticaily as iron content

increased. At first it would appear that increases in iron graduaiiy change the ovedi

chemistry of the alloy enough to cause it to solidify as a hypoeutectic aiioy. Presumably,

when AisFeSi crystaliizes as the first phase to grow, siiicon is absorbed into its structure,

depieting the melt sufficiently to shift from near eutectic composition to a hypoeutectic

composition. This behaviour would correlate weïî with the fact that the eutectic duration

decreases with increasing iron, and the duration and formation temperatures of the a-

aluminum heat e ffect increase as iron increases.

An explanation for these observations can be made using the sirnpiified Al-Fe-Si-

Mn phase diagrarns used before by other re~earchers'.~~ to explain the growth of Fe-

bearing phases in alurninurn-silicon ailoys. Figure 6.4 shows a sirnplified quaternary Al-

Fe-Si-Mn phase diagram with constant 0.25% wt Mn (the Mn level for the 413 dloy

studied). At low iron, both Al15(Mn,Fe)$i2 and AisFeSi were observed in the

microstructure. This agrees with the solidification directions seen within the phase

diagram. The chinese script f m s first, and the composition moves to the A1peSi and

Ai15(Mn,Fe)Si2 valley where both iron bearing intermetallics form. Eventually the

chemistry changes due to segregation and the Ai-Si-/3 temary eutectic begins to fonn. At

higher iron contents the first phase to form is the A4FeSi intermetallic. Once segregation

occurs due to AlPeSi growth, the overaii composition wili be in the

CHAITER SIX : DISCUSSION OF RESULTS 94

Simplified Al-Fe-Si (Constant Mn) 0.25% wt Manganese

Siliocn %wt

Figure 6.4 : The two data points in this figure represent the siliocn and iron contents of two of the thermal analvsis sam~les studied.

CHAPTER SIX : DISCUSSION OF RESULTS 95

A15FeSi/Ai15(Mn,Fe),Si2 valley. Here both A15FeSi and Al15(Mn,Fe)& crystals begin to

form, and the AlPeSi phases are much longer while the Al,@ln, Fe)& phases are more

star-like. Findy the chemistry of the remaining iiquid achieves the ternary eutectic.

It shouid be noted that, unlike the pre-eutectic A15FeSi phase observed in the

319.2 and A356 alioys, the A1,FeSi in high iron A413 is predendritic. It does not form

in the presence of aluminum dendrites. As a consequence, the Al,FeSi that formed at

high iron contents in 413 was not cooling near equilibriurn cooling rates due to the

absence of the aluminum latent heat effect. At equiiibrium rates the AisFeSi should have

been resolvable at the base 0.90% wt iron for the 413 aUoy, according to the Al-Fe-Si

temary phase diagram shown in chapter two. In effect the 4FeSi thermal anomaly in

413 aüoy was resolvable on cooling curves only at 1.9% wt iron.

The variation in eutectic duration with increasing iron content, due to the

depletion of silicon from the melt, is more dramatic and varies for the iron range studied.

This is opposite to the case for the hypoeutectic alloys studied. In 4 13 alloy , eutectic time

duration is probably the best parameter to correlate with irOn concentration.

6.1 Quality of Derivative Curves

The l d maxima of the dwde cuwes used for the determination of start and end of

solidification times were used to measure the most important parameters associated with

the A1,FeSi intermetallic and aluminum-siliwn eutectic. However it is evident on these

denvative curves that other peaks of lesser magnitude occur and sometimes merge with

the local maxima. The ASCII file exported from the Research Meltlabm Software shows

that the temperature value measured at each sampling has a slight fiactional difference

from the previous data point. This is believed not to be an electronic problem, but rather

a problem uising from the sarnple setup. The thermal analysis setup to achieve very slow

cooling rates (insulated cup method) allows heat to be extracteci mainly through the top

of the sample while the sides and bottom are insulated with f i b e h . The directionality

of the heat extraction coupled with the slow cooiing rates give rise to srnail convection

cells during solidification. The thermocouple would then expenence a slight cyclic

temperature change superimposeù on the gradua1 decrease in temperature of the whole

C H A P T E R ~

sample. This could result in fractional differences in temperature readings fiom one point

to the next.

A second problem experienced in this work was that sample size had an effect on

the repeatability of the time parameters associated with certain thermal anomalies.

Quantitative thermal analysis obviously requires excellent control of sarnple mass.

CHAPTER SEVEN : CONCLUSIONS & FUTURE WORK 97

Chapter Seven

@onclusions & Future Work

7.0 Conclusions Themal anaiysis of three aiuminum-silicon alloys having different iron contents was

studied for this thesis. The objective was to establish the possibility of using thermal

analysis to quanti@ the iron content of an alurninum-silicon melt prior to the casting

operation. The main conclusions that can be drawn from the thesis are :

1 .) Maxima on the d?/dt2 cuwe c m be used to identify the initial formation time and

temperatures of the A15FeSi intermetallic, the aiuminum-siliwn eutectic, Mg2Si-Al

eutectic and CuAl,-Al eutectic.

2.) As the iron content is increased, the resolvable formation temperature of the

A15FeSi phase, as signalled by a local maximum on the dq/dt2 curve, increa~e~

for the three alloys studied.

3.) For the two 300 series alloys studied (319.2 & A356) the A1,FeSi themal

signature was observable within the region of the cwling cuve between the

prirnary arrest and the beginning of the eutectic plateau. The minimum iron

contents detectable were 0 . 5 4 % ~ (356) and 0.64% wt (319) using the insulated

cup method (O.lCPC/sec). Both the time duration and resolvable formation

temperature of the A1,FeSi intennetallic were measured, and it was shown that

both could be reliably used to quanti@ the iron content in the melt.

CX.AFI'ER SEVEN : CONCLUSIONS & FUTURE WORK 98

4.) In A356 alloy mled at O.IO"C/sec the measured formation temperatures of the

A15FeSi intermetallic plotted on an Al-Fe pseudo-binary lay on the liquidus line

of the Ai+@+liquid field. At 0.4S°C/sec this data was shifted to the right with

respect to the pseudo-binary phase diagram, resulting in an adjustment of the

range of iron detectabiiity from 0.54% wt-1.3 % wt to 0.75 % wt- 1.45 % wt.

5. ) In 319.2 aüoy the formation temperature and time duration of the A1-i was

repeatable for trials one and two, giving confidence that the thermal analysis

method can be used reliably to determine the iron concentration of the melt.

6.) For the 400 series alloy studied (A413), the A1,FeSi thermal signature was

observed prior to pnmary a-aiuminum and alurninum-silicon eutectic only at iron

contents above 1.9% wt. It was found that a more dramatic change in eutectic

time duration result for the whole iron range studied.

6.) Sarnple mass is a critical variable to control in thermal analysis if tirne based

parameters are to be used.

CHWTER SEVEN : CONCLUSTONS & FUTURE WORK

7.1 Future Work The objective of this thesis project was to establish whether or not it could be possible

to quantxfy the iron content of aluminum-silicon aüoy melts within a reasonable range

prior to the casting process. It was determined that for hypoeutectic aüoys this was

possible for 356 aUoy and 319 aiioy. For the 413 aUoy studied, iron within the standard

ranges was detected only indirectly through the effect on the alurninum-silicon eutectic

duration. The feasibility study perforrned on these three alloys could be expanded to

other casting and wrought Uoys where iron is a signifiant element in the chemistry.

Use of a time parameter for eutectic modification determination needs further

study. It has been postulated that the eutectic time increases with modification. The use

of the local maximum in the d2~/dt2 cume for hypoeutectic alloys, and the use of maxima

or minima on the d2T/dt2 curve for near eutectic aiioys could be used to mathematically

determine the start and end times for eutectic growth, and thus infer on the degree of

modification. This method of detennining time for the modification of the aiioy however

relies on the sample size being set and controiïed to a firm standard.

REFERENCES KI0

References

1. Gilbert F. Bourcier, John Dickinson, Joseph Tessandon & David Schiffer, Aluminum Recycling Casebook, The Alurninum Association, Dec. 1985,

2. John E. Gnuleski & Bernard M. Closset, The Treatment of Liquid Alurninum-Silicon Aiioys, Amencan Foundrymen' s Society, Inc., 1990, 256pp.

3. Aluminum Casting Technology, Revised by Jorstad, J. L., Rasmusen, W. M., Edited by Zalensas, D. L., 2nd Ed., Amencan Foundrymen's Society, Des Plaines, Ili, 1993.

4. Q. S. Harned, M. Dogan & R. Elliot, " The Dependence of Secondary Dendrite Arm Spacing on Solidification Conditions-III. Al-7Si-O.SMg Alloys Treated with TiBAl and TiBAl/Sr Additions", Cast Metals, Vol. 6, Nurnber 6, 1993, pp.47-53

5. F. Paray & J. E. Gruzleski, "Microstnicture-Mechanid Property Relationship in a 356 Alloy. Part 1: Microstructurew, Cast Metals, Vol. 7, number 1, 1994, pp. 29-40

6. F. Paray & J. E. Gnizleski, "Microstnicture-Mechanical Property Relationship in a 356 &y. Part 2: Mechanical Properties", Cast Metals, Vol. 7, number 3, 1994, pp. f 53-163

7. F. Paray & J. E. Gruzleski, "Factors to Consider in Modificationw, AFS Transactions, Vol. 102, 1994, pp.833-842

8. L. Backenid, G. Chai, J. Tamminen , "Solidification Characteristics of Aluminum AUoys" , Vol. 2, Foundry Alioys, AFS/Shanaluminum, 1990, pp.255

9. F. Paray & J. E. Gruzleski, "Modification- A Parameter to Consider in the Heat Treatment of Al-Si Alloys", Cast Metals, Vol. 5, Number 5, 1993, pp. 187-197

10. G. Gustafsson, T. "Thorvaldsson, & G. L. Dunlop, " The Influence of Fe and Cr on the Microstructure of Cast Al-Si-Mg Alloys" , Metallurgical Transactions A, Vol. 17A, Jan. 1986, pp. 45-52

11. P. N. Crepeau, " Effect of Iron in Al-Si Alloys: A Critical Review", AFS Transactions, Vol. 103, (1995), pp. 361-366

12. K. H. Kirgin, " 1996 Slowdown Sets Stage For 1997 Resurgencen , Modem Casting, AFS, January 1996, pp.27-30

References 101

13. M. L. Philbin,"Diecasting 1996 Status Reportw, Modern casting, AFS, February 1996, pp.30-33

14. Joseph C. Runkle," Tool-Steel Composites Offer Improved Wear, Fatigue", Advanced Materials & Processing, Vol. 150, Number 2, Aug. 1996, pp.7

15. S. G. Shabestari & J. E. GNzleski, " Gravity Segregation of Complex Intermetailic Compounds in Liquid Aluminum-Silicon Ailoysn, Metallurgical Transactions A, Vol. 26A, April 1995, pp. 999-1006

16. J Lanker, " Non-Equilibnum Solubility of Iron In Aluminum " , Key Engineering Materials, Vol 44 & 45 (1990) pp. 135-146.

17. H. W. L. Phillips," Annotated Equilibrium Diagrams of Some Aiuminum Ailoy Systems", Institute of Metais Monograph & Report Series, (1959). pp. 56-65

18. L. F. Mondolfo,"Aluminum Alloys : Structure and Properties", Butterworths, 1976, pp. 759-777, pp. 534-37, pp. 661-62

19. V. G. Rivlin and G. V. Raynor, " Criticai Evaluation of Constitution of Aluminum- Iron-Silicon Systemn , International Metals Reviews, No. 3, 198 1, pp. 133-35

20. T. Turmezey, V. Stefaniay and A. Gnger,"ALFeSi Phase in Aluminumn, Key Engineering Materiais, Vols. 44 & 45, Trans Tech hiblications, Switzerland, 1990, pp. 57-68

2 1. J. Iglessis, C. Frantz, M. Gantois, "La Croissance par Propagation de Plans de Macle des Composes Eutectique dans le S ysteme Ternaire Al-Fe-Si " , Memoires Scientifiques Revue Metallurgie, Feb. 1978, pp. 90-10

22. R. Hoier, O. Lohne, S. Mortvedt,'AEeSi-Particles in an Al-Mg-Si-Fe AUoyn ,Scandinavian Journal of Metallurgy, No. 6, Vol. 1, 1977, pp.36-37

23. P. Skjerpe, "Intermetallic Phases Formed during DC-Casting of an Al-0.25Wt %Fe- 0.13Wt%Si Alloyw, Metall. Trans A, Vol. 18A Feb. 1987, pp. 189-200

24. M. H. Mulazirnoglu, A. Zaiuska, J. E. Gruzieski & F. Paray , " Electron Microscope Study of Al-Fe-Si Intermetaliics in 6XXX Senes Aluminum Alloysn, Metallurgical and Materials Transactions A, Vol. 27A, April 1996, pp. 929-936.

25. M. H. Muiazimoglu, N. Tenekedjiev, B. M. Closset, J. E. Grezleski, "Microstructure and Thermal Analysis of Strontium Treated Aluminurn-Silicons AUoys" , Amencan Foundryrnen's Society Inc., (1995), 1 l6p.

eferences 102

26. 0. Vorren, J. E. Evensen & T. B. Pedersen, ' Microstructure and Mechanical Properties of AlSi(Mg) Casting Alloys" AFS Transactions, Vol 93, (1984) pp. 459466

27. S. Shivkumar, L Wang & D. Apelian, " Molten Metal Processing of Advanced Cast Aluminum AUoys" , AFS Transactions, Vol 97, (l989), pp. 8 1 1-824

28. L. Ananthanarayanan, F. H. Samuel, J. E. Gruzleski, " Thermal Analysis Studies on the Effect of Cooling Rate on the Microstnicture of 319 Aluminum Alloy", AFS Transactions, (1992), Vol. 99, pp. 383-391

29. W. Bonsack, " Iron-The Problematic Factor in Quality of Aluminum AUoy Die castings" AFS Transactions, (l961), Vol. 69, pp.712-720

30 S. G. Shabestari & J. B. Gruzieski, "Modification of Iron Intermetaliics by Strontium in 413 AUoysn, AFS Transactions, (1995), Vol 103, pp. 285-293

3 1. B. Xiufang, C. Guohua & M. Jiaji, " A Master Alloy for the Speroidisation of Needle-Form Iron Compounds in Aluminum AUoys" Cast Metals, Vol. 6, # 3, 1993, pp. 159-161

32. B. Xiufang, 2. Guohua, 2. Shengxu & M Jiaji, " The Speroidisation of Needle-forrn Iron Compounds in an Ai-Si AUoy", Cast Metals, Vol. 5, # 1, 1992, pp.39-42

33. L. Ananthanarayanan, F. H. Samuel, & J. E. Gruzleski, " Crystallization Behaviour of Iron-Con taining Intermetallic Compounds " , Metallurgical & Materials Transactions A, Vol. 25A, Aug. 1994, pp.1761-1773

34 J. Gauthier & F. H. Sarnuai, "Tensile Properties & Fracture Behaviour of Solution - treated 3 19.2 Alloy Al Automotive Alloys", AFS transactions, (1995), Vol. 103, pp. 849-857

35. N. Crowell & S. Shivkumar, "Solution Treatment Effects in Cast Al-Si-Cu Aiioysn, AFS Transaction, (1995). Vol. 103, pp. 721-726

36. N. Tenekedjiev, J. E. Gruzleslci, ' Thermal Analysis of Strontium treated Hypoeutectic and Eutectic Alurninum-Silicon Alloys" , AFS Transactions, ( 199 1). Vol. 99, pp. 1-6

37. M. H. Mulazimoglu, N. Tenekedjiev, B. M. Closset, J. E. Gruzleski, " Studies of the Minor Reactions and Phases in Strontium-Treated Aluminum-Silicon Casting Aiioys", Cast Metals, Vol. 6, # 1, 1993, pp. 16-28

38. S. Gowri, " Cornparison of Thermal Analysis Parameters of 356 and 139 Alloys" AFS Transactions, Vol. 102, 1994 , pp. 503-508

39. D. Sparkman, A. Keamey , " Breakthrough in Aluminum Aiioy Thermal Analysis Technology for Process Control", AFT Transactions, Vol. 102, 1994, pp. 455-460

40. D. Sparkman, "Understanding Thermal Analysis of Iron" , AFS Transactions, (l994), Vol. lm, pp. 229-233

41. F. J. Bradley, P. F. Bartelt, C. A. Fung & R. W. Heine, "Considerations in the Applications of Numencai Simulation to S hnnkage Predictions in Ductile Iron castings " , AFS Transactions, 1992, Vol. 100, pp. 155-162

42. Editiors: J. R. Davis " ASM Specialty Handbook: Aluminum & Aluminum AUoysn,

IMAGE EVALUATION TEST TARGET (QA-3)

APPLIED - IINLAGE . lnc = 1653 East Main Street - -. - Rochester. NY 14609 USA =-A Phone: i l 6/482-0300 Fa: 7161288-5989

quantification of iron in alsi foundry...

Documents