r

The Rol~ of Magnesium

on

Eutectic Silicon Microstructure

By

Ahmad. T. Joenoes

A Thesis Submitted to the

Faculty i)f Graduate Studies and Research

in Partial Fulfilment of the Requirements for the Degree of

Master of Engineering

Dcpal1ment of Mining and Metallurgical Engineel'ing

McGiII University

MONTREAL - CANADA

March 1991

i

ABSTRACT

To improve the mechanical properties of the most widely

used Al-Si foundry alloys, small amounts of magnesium are

often intr-oduced. However, the effect of magnesium on the

eutectic microstructure is unclear. The aim of the present

work is to investigate the effect cf magnesium on the eutectic

microstructure and te clarify its role in the silicon

solidification.

Experiments were performed on hypoeutectic and eutectic

alloys (Al-7%Si, Al-ll%Si, Al-13'Si and 413.2 aJloys).

Assessment of the eutectic microstructun! was don€: using

thermal analysis and metallography. The silicon phase was

quantified by image analysis in terms of its perimeter, shape

factor and aspect ratio. The results reveal that in hypo­

eutectic and eutectic synthetic alloys magnesium decreases the

homogeneity of the micrQstructure. At about l%Wt Mg, however,

mag:1esium refines slightly the silicon phase. It was found

that magnesium has a negative effect on strontium modifi­

cat ... on, that is, it changes the microstructure from weIl

modified to partially modified. This is believed due to the

presence of a complex intermetallic compound (Mg2srA14si3 )

formed prior to the eutectic transformation. It is suggested

that magnesium not be added in excess of 0.6 Wt% in strontium

TIlodified Al-si alloys.

..

. '

ii Résumé

Pour améliorer les propriétés roécaniques des alliages de

founderie Al-si les plus largement utilisés, de petites

quantités de magnésium sont souvent ajoutées. Cepend.:'nt,

l'effet du magnésium sur la structure eutectique n'est pi1S

clair. Le but de ce travail a été de déterminer l'cff0t

du magnésium sur la structure eutectique ~t de clarifier son

rôle dans la solidification du silicium.

Pour cette étude, des alliages hypoeutectiques et

eutectique ont été utili~és (alliages 413.2, Al-7%Si, Al-l1%Si

et AI-13%Si). L'évaludtion de la structure eutectique a été

faite pdr analyse thermique et par métallographie. La

morphologie du silicium ~ été quantifiée par analyse d'image

(mesures de la taille et de la forme). Les résultats montrent

que pour les all iages synthétiques (eutectique et hypo­

eutectiques), le magnésium diminue l'homogénéité de la

microstructure. Cependant, la présence de 1% en poids de

magnésium affine légèrement la phase de silicium. Il a été

trouvÉ que le magnésium a un effet négatif sur la modification

par le strontium. En effet, la structure qui était bien

modifiée, devient partiellement modifiée. Ceci s'expliquerait

par la présence d'un composé i~termétallique complexe

Mg2SrAl4Si3 , formé avant la transformatioll eutectique. On

suggère de ne pas ajouter plus de 0.6% en poids de magnésium

aux alliages Al-5i modifiés ~u strontium .

1 ..

l would

gratitude to

iii

AC~NOWLEDGEMENTS

U ke to express my sincere appreciation and

Prof essor J.E. Gruzleski for his excellent

supervision, const3nt and constructive encouragement and

patience throughout the course of this research.

Special thanks are due to N. Tenekedjiev, F. Paray, C.

Aliravci and P. Blanford for their valuable suggestions and

discussion during conducting this work.

The technical assistance of R. Salby and N. Handiak for

casting experiments, H. Campbell for electron microscope

works, and M. Knoepfel for mechanical works i5 aiso gratefully

acknowledged.

Finally, l am also indebted my parents, colleagues and

friends for their generous support and contribution to this

work in various ways during the completion of this research

and my study at McGill University.

.,

iv Table of Contents

Page

Abstract Acknowledgem~nts

Table of Contents

List of Figures

List of Tables

• • • • • • • .. .. .. .. .. .. • • • • • • • .. • • • • • • • .. • • • •• 1

• .. • • • .. • • • • • • • . .. • .. .. .. .. .. .. .. .. .. .. .. .. .. .. .... 111

.. .. .. .. . . .. .. .. .. .. .. .. . .. .. .. .. .. .. .. .. .. .. .. .. ... . .. .. . .. i v

.. .. .. .. • • .. .. .. .. . .. .. .. .. .. .. .. .. .. .. .. .. .. . .. .. .. .. .. .. ... vi

.. .. .. . . . .. .. .. .. . . .. .. .. .. .. .. .. .. .. .. .. .. .. .. .. .. .. . .. ... i x

Chapter l INTRODUCTION • .. .. .. .. .. .. .. .. .. • .. .. .. .. .. .. .. .. .. .. .. .. .. .. .. .... 1

1.1 - The Market for Aluminum

1.2 1.3 1.4

1.5 1.6

1.7

1.8

casting Alloys .•.•.....•........... 1 - Al-si castings Alloys . . . . . . . . . . . . .. 3 - Typical Examples of Al-si Alloys 5 - Mechanical Properties and Micro-

structure of Al-Si Alloys .......... 10 - Modification •....•................ 14 - The Role of Thermal Analysis

in Modification Assessment .....•... 21 - Quantitative Analysis an~ Particles

Character i zation . . . . . . . . . . . . . . . .. 24 - The Aims of The Present Work ..••.. 26

Chapter II EXPERIMENTAL PROCEDURE .................. 28

2 .. 1 - General ... . .. .. . .. .. .. .. .. .. .. .. .. .. .. .. .. . .. .. .. .. .. .. ... 28 2.2 - Alloy Preparation & Alloy Making •... 28

2.2.1 Materials .•.•....•......••... 28 2.2.2 Melting & castings ........•... 30

2.3 - Thermal Analysü~ ..•.............•... 3:2 2.4 - Spectrochemical Analysis .......•... 34 2.5 - Metallographie Examination .. .. .. .. .. .. .. .... 34

2.5.1 Optical Microscopy 34 2.5.2 Scanning Electron Microscopy .. 36

2.6 - Quantification Using Image Analysis ..................... 36

(

Chapter III RESULTS & DISCUSSION • . . . . . . . . • • . . . . . . .. 43

3.1 - Magnesium & strontium Effect in Hypoeutectic synthetic Alloy ....... 43

3.2 - Magnesium & strontium Effect in Eutectic Synthetic Alloy ...•........ 52

3.3 - Magnesium & strontium Effect in 413.2 commercial Alloy ...•.•.......• 62

3.4 - Magnesium & strontium Effect on the Microstructure of 356.2 and 357.2 Commercial Alloy .........•.•.•...... 67

3.5 - Effect of Magnesium on th~ Eutectic Temperature ..•............. 68

3.6 - Image Analysis Results .............. 74 3.7 - Aspects to be Considered ............ 83

Chapter IV CONCLUSIONS . . • • • • • . • • . . . . • . . . • . . . . . . . . .. 89

References ............... ,. • • • • • • • .. • • • • .. .. .. • • .. .. .. .. .. .. .... 90

Appendices • • .. .. • • • • .. .. • .. • • • • • • • • • • • • .. .. .. • .. .. .. • .. .. • .. • .. •• 97

..

VI List of Figures

Figure 1.1 Aluminum-silicon phase diagram showing composition range of various commercial alloys

Figure].2 Relationship between mechanical properties and the eutectic microstructure 1 ~

Figure 1.3 The alteration of the cooling curve due ta modification , , . )

Figure 1.4 Cooling curve and its corresponding derivative curves for Al-7%Si-1Mg synthetic allay 2.1

Figure 1.5 The variation of silicon particle diameter and aspect ratio with strontium concentration 26

Figure 2.1 Experimental Procedure il 'J

Figure 2.2 Casting & Thermal Analysis System J î

Figure 2.3 Molten metal of Al-si alloy where magnesium

Figure 2.4

Figure 2.5

Figure 2.6

Figure 2.7

Figure 3.1

Figure 3.2

Figure 3.3

Figure 3.4

and strontium were introduced 3J

Sample Preparation 35

Scanning electron microscope coupled with an image analysis system 37

Particles Characterization 38

Image Analysis Block Diagram 41

The cooling curve for unmodified Al-7%Si alloys 46

The cooling curve for modified AI-7%Si alloys 47

A typical cooling curve for Al-si alloys 49

The evolution of eutectic microstructure as a function of magnesium in unmodified Al-7%si synthetic alloy 50

Figure 3.5 The evolution of eutectic microstructure as a fUt"cticl1 of magnesium in strontium

vii

modified Al-7\si synthetic alloy 51

Figure 3.6

Figure 3.7

The evolution oi eulectic microstructure as a function of rnagnesium in unmodified Al-11\Si synthetic alloy

The evolution of eutectic microstructure as a function of magnesium in strontium

54

modified Al-11%Si synthetic alloy 55

Figure 3.8

Figure 3.9

The cooling curve for unmodified Al-13%Si alloy

The cooling curve for strontium modified Al-13\si alloy

Figure 3.10 The evolution of eutectic microstructure as a function of magnesium in unmodified Al-13%si synthetic alloy

Figure 3.11 The evolution of eutectic microstructure as a function of magnesium in strontium

57

58

60

modified Al-13%Si synthetic alloy 61

Figure 3.12 The evolution of eutectic microstructure as a function of magnesium in unmodified modified 413.2 commercial alloy 65

Figure 3.13 The evolution of eutectic microstructure as a function of magnesium in strontium modified 413.2 commercial alloy 66

Figure 3.14 The eutectic microstructure of a modified 356.2 & 357.2 alloy

Figure 3.15 The depression of the eutectic temperature as a function of magnesium content for un-

67

! modified & strontium modified Al-si alloys 69

Figure 3.16 Part of a ternary Al-Si-Mg phase diagram 71

.,

V 11 i

Figure 3.17 Secondary Electron Image

Figure 3.18 Binary images which are created from

video/secondary electron image

Figure 3.19 The perimeter of silicon particles as a

function of magnesium concentration

75

76

(Al-7%Si alloys) 19

Figure 3.20 The shape factor of silicon particles as a

function of maqnesium concentration

(Al-7%Si alloys) 79

Figure 3.21 The aspect ratio of silicon particles as a

function of magnesium concentration

(lU-7%Si alloys)

Figure 3.22 The perimeter of silicon particles as a

function of magnesium concentration

80

(Al-13%Si alloys) 31

Figure 3.23 The shape factor of silicon particles as a

function of magnesium concentration

(Al-13%Si alloys)

Figure 3.24 The aspect ratio of silicon particles as a

function of magnesium concentration

82

(Al-13%Si alloys) 82

Figure 3.25 Intermetallic compound found in Al-Si-Mg-Sr

system

Figure 3.26 The complex Mg2SrAl4Si} intermetallic

compound

35

86

.

List of Tables

Table 1.1 U.S. Aluminurn Casting Shipments

Table 1.2 composition of the most common

a lumin 1 lTl1-si l icon casting alloys

Table 1.3 casting characteristics of alurninum­

silicon alloys

Table 1.4 Typical mechanical properties of

aluminum-silicon alloys

Table 3.1 The effect of mdgnesium on Al-7%Si

ix

Page

2

6

6

10

synthetic alloys 48

Table 3.2 The effect of magnesium in the presence

of strontium on Al-7%Si synthetic alloys 48

Table 3.3 The effect of rnagnesium on Al-11%Si

synthetic alloys 53

Table 3.4 The effect of magnesium in the presence

of strontium on Al-l1%Si synthetic alloys 53

Table 3.5 The effect of magnesium on Al-13%Si

synthetic alloys 59

Table 3.6 The effect of rnagnesium in the presence

of strontium on Al-13%Si synthetic alloys 59

Table 3.7 The effect of magnesium on 413.2

commercial alloys 64

Table 3.8 The effect of rnagnesium on 413.2 in the

presence of strontium commercial alloys 64

Table 3.9 Linear relationship between the eutectic

temperature and the magnesium level 68

Table 3.10 Comparison of the eutectic temperature

obta i ned from exper iment and those from

ternary phase diagram 73

'f !

4

1

Chapter 1

1 N T R 0 DOC T ION

1.1 - The Market for Aluminum casting Alloys.

Aluminum, since the beginning of the middle of this

century, is perhaps the most fascinating light metal used for

large scale applications. The excellent results achieved by

alloying with other elements lead to an almost unlimitcd

variety of properties. The contribution of aluminum alloys as

cast materials, in terms of tonnage, in 1986 is second largest

only to ferrous castings l• As an illustration, in 1986, US

shipments of this alloy were above one million tons. Aluminum

cast parts represented a~proximately 10% of the total tonnage

of US casting shipments. The current market for aluminum

castings is supplied by the most commonly used processes which

are sand casting, permanent mold, and high pressure die

casting2• Table 1.1 shows the aluminum casting shipments (1000

tons per year) in the U.S. alone. The table indicates that

during the period between 1945-1986 the shipments of aluminum

casting alloys have increased by almost ten fold3•

Among others, the growth of high pressure die casting is

the most substantial, with an increase from 34,000 tons in

1945 to 825, 000 tons in 1986. This number will probably

increase in the following decade. The significant rise in die

castings is due to the fact that this process produces metal

2

sections thinner than those obtainable by other techniques.

High mechanical properties and resistance to fatigue failure

can be achieved without heat treatment because of the fine

microstructure produced naturally by rapid solidification4•

Table 1.1 U.S. Aluminum Casting Shipments (1000 Tons Year) [3]

Sand casting Permanent Mold High Pressure Year casting Die Casting

Tons % Tons % Tons %

1945 98 52.5 54 29.0 34 l.8.5

1955 83 20.0 149 36.5 178 43.5

1963 72 15.0 150 32.0 254 53.0

1983 83 9.4 151 17.0 653 73.6

1986 85 7.8 170 15.7 825 76.4

Total

Tons

186

410

476

887

1080

The rapid expansion of aluminum casting by high pressure

die casting cannot be separated from the revolution in the

automotive industry. The rapid economic growth in the world l

and the sharp oil price increases in the early 1970's resulted

in a trend to smaller,lighter and more economical automobiles.

The energy crisis during the past ten years in most developed

countries has led to a search for materials that can increase

the efficiency of ail usage. Aluminum alloy castings provide

cost-effective answers to many needs and problems in the

automotive industry. The combination of good mechanical

properties and casting characteristics result in a variety of

-.

3

applications including engine blocks, cylinder heads, intakc

manifolds, pistons, wheels, etc 3J.

Another market segment for aluminum casting alloys is in

both the civilian and military aircraft industries. In

addition, major uses for aluminum castings are found in

computers, motors, generators and communication systems. In

~he U.S alone, the market opportunities for aluminum casting

alloys over the next decade are expected to grow as the

demand for all kinds of castings rises. To fulfil this high

demand, it follows that, not only must novel techniques be

employed in casting processes, but it is also necessary

to have improvements in the level of the cast aluminum

properties. This means the need for better mechanicùl

properties, through microstructural control, at the lowest

production cast.

I.2 - Al-si casting Alloys.

Alloys with silicon aS the major alloying ingredient are

the most important commercial casting alloys, prirnarily

because of their superior casting characteristics in

comparison to other aluminum alloys. Additions of silicon to

pure aluminum dramatically improve fluidity, hot tear

resistance, and feeding characteristics3 • other advantages

of these alloys are high resistance to corrosion,_ good

weldability and the fact that silicon reduces the coefficient

.. ,

4

of thermal expansions,7. The aluminum-silicon castings alloys,

in fact, constitute 85% to 90% of the total aluminum cast

parts produced. These commercial alloys are available with

hypoeutectic (5%-1 O%Si) , eutectic (11 %-13%8i) and less

commonly hypereutectic (14%-20%8i) compositions6 • Figure 1.1.

shows the binary aluminum silicon phase diagram and the range

of these three different compositions. The choice of alloy is

on the basis of the desired mechanical properties required by

the particular application. The higher the silicon content,

the harder and stronger the materlal, but at the expense of

ductili t y5, 7,8,9. Generally, the optimum range of silicon ~s

assigned to casting processes based on the relationship

between cooling rate, fluidity, and the amount of eutectic

phase on feeding8 •

Aluminum casting alloys are classified on the basis of

their chemical composition. This classification varies from

one country to another. In the U.S., the Aluminum Association

has adopted a four-digit numerical system to identify aluminum

casting alloys'o. The aluminum silicon alloy family is

designated as folllJws 3xx. x for silicon wi th magnesium

and/or copper as a major impurity (typically >3%) and 4xx.x

for silicon w~thout appreeiable chemieal impurities. The last

two digits have no specifie signifieance to identify the

diff~rent alloys in the system. The last digit to the right

5

of the decimal point identifies the product form either as an

ingot (0) or as a final cast product(l) .

Oc

100

600

500

o AI

Llquld

5 10 15

Wt%

u :0:

-::1 L+SI Q) ..

Q) 0->0-Z

20 SI

Figure 1.1 Aluminum-Silicon phase diagram showing composition range of various commercial alloys (6].

Table 1.2 indicates the most common aluminum silicon casting

alloys with their various chemical compositions. The relative

casting characteristics achieved through these compositions

are listed in Table 1.3.

1.3 - Typical Example. of Aluminum-Silicon Alloys.

Hypereutectic alloys with a high silicon content are

characterized by outstanding wear resistance, a lower thermal

expansion, and have the best casting characteristics ll •

6

1 Table 1.2 Composition of the most common aluminum-silicon "-

casting alloys [3].

Alloy Method*) Elements (wt%) Si Cu Mg Fe Zn Others

319.0 S,P 6.0 3.5 <0.10 <1. 0 <1.0 332.0 P 9.5 3.0 1.0 1.2 1.0 355.0 S,P 5.0 1.25 0.5 <0.06 <0.35 A356.0 d,P 7.0 <0.20 0.35 <0.2 <0.1 A3~7.0 S,P 7.0 <0.20 0.55 <0.2 <0.1 0.05Be 380.0 0 8.5 3.5 <0.1 <1.3 <3.0 383.0 0 10.0 2.5 0.10 1.3 3.0 O.15Sn 384.0 0 11.0 2.0 <0.3 <1.3 <3.0 0.35Sn 390.0 0 17.0 4.5 0.55 <1.3 <0.1 <O.lMg 413.0 0 12.0 <0.1 <0.1 <2.0 443.0 S,P 5.25 <0.3 <0.05 <0.8 <0.5

,~ *) S, Sand Casting; P, Permanent Mold Casting t \. D, High Pressure oie Casting.

Table 1.3 Casting characteristics of aluminum-silicon alloys [ 3 ] •

Alloy casting Resistance Pressure Fluidity Shrinkage method ta tearing tightness tendency

319.0 S,P 2 2 2 2

332.0 P 1 2 1 2 355.0 S,P 1 1 l 1 A356.0 S,P 1 1 l 1

A357.0 S,P 1 1 l 1 380.0 0 2 2 2

390.0 0 2 2 2

413.0 0 1 1 l

443.0 P 1 1 2 1

, 1 Ratings 1, best . 5, worst. ,

7

However, those alloys have shortcomings, due to the presence

of the extremely hard primary silicon phase that reduces

considerably the life of tools during machining. By contrast,

the other two types of aluminum silicon alloys, eutectic and

hypoeutectic, are attractive due to the wide range of the

properties offered. Hypoeutectic alloys, for example alloy 443

(5.3% silicon), can be used with aIl casting processes in

which good ductility, good corrosion resistance and pressure

tightness are more important than strength8•

Chemical impurlties strongly affect the final mechanical

properties. Iron sometimes is introduced in small amounts to

the hypoeutectic alloy to irnprove hot tear resistance and to

decrease the tendency of die sticking or soldering in die

casting. Increasing the iron content (>0.8%), however, is

reported to decrease the ductility by forming hard brittle

plates ~f Q-AlFeS~~. The detrimental effect of Iron on the

mechanical properties can be alleviated by using heat treat­

ment. Alternatively, Co,Cr,Mn, Mo, and Ni are sometimes added

to correct for iron7•

Alloy A444 (7%Si-O.2%Fe max.) exhibits good corrosion

resistance and especially high ductility when it i5 cast in

a permanent mold, followed by heat treatment to the T4

condition8 • This alloy is used when impact resistance is of

primary concerne Strengthening of the hypoeutectic alloy can

be achieved by introducing small amounts of Cu, Mg, or Ni.

Copper (4% max.) imparts moderately high strength and improved

8

rnachinability. A slightly higher content of this elernent is

associated with reduced ductility and lowered corrosion

resistance7• Alloy 319 (6%5i-3. S%Cu) is used for sand and

permanent mold castings. Alloy 380 (8.5%Si-3.5%cu) is

preferred for die casting, while 333 alloy (9%Si-0.5%CU) was

developed for permanent mold casting. Many castings of this

general-purpose alloy are supplied only in the as-cast temper,

but the strength can often be improved through full T6 or T7

heat treatment8•

Another group of aluminum silicon alloys are those that

conta in magnesium. The most popular of these alloys is 356

(7%5i and O. 3%Mg) or the higher purity A356. Alloy 357 is

similar to 356, but contains a higher magnesium level (0.5%).

The alloys are marked by excellent casting characteristics,

weldability, pressure tightness and corrOSlon resistance8 •

These two alloys are heat treatable due to M92Si (magnesiurn

silicide) hardening. The hard particles of Mg2Si are

precipitated uniformly throughout the aluminum matrix after

solution treatment (T4), followed directly by quenching and

aging8 • Heat treatment provides the variety of tensile

s~rengths and other physical properties that are attractive

for many applications, including military, aircraft and

automotive parts. The hardening effect, however, is useful

only to approximately 0.70% Mg, beyond which point no further

9

strengthening in the aluminum matrix takes place ann the 10ss

of ductility is even more pronounced12 •

It is weIl known that prolonged or repeated heating will

cause silicon to spheroidise, and as a result, it irnproves

impact resistance IJ•14 • However, in some less critica1 casting

applications, such as in permanent mold or die-cast parts,

solution treatment and quenching May be eliminated. Rapid

cooling in the rnold is relied upon to produce strengthl~.

There are Many other alloys with different chemica1

constituents and casting processes that were not rnentioned

here. Table 1.4 lists some of the Most common aluminum silicon

alloys associated with their casting processes and the

resulting mechanical properties.

, '~ ~'"

'1'IbIe i.~ TnDIIIe"tw"œI Pqertie& ~ Aluminum-gJbm ADD.7I rn

. SjiI'Oii'iIIt _ft(.. . ..... ... ·yt.. . C'JS

.. l\ii ... ...-t JJIé:lilt .... '. f:":" ~.~':"·1· . ::: ..... :. .f ~ ." .' . 1ÎIP!I!i' ••• :.1IIitd:' ... ';,-..' .. .... .-,.... . DIT'" ... . .. 1l!S... ... ..iWIi-.... . •• ., . . . ·Utll-. ". '. 7l6.. ' .

AI-!I:5-7m n.m 400-500 80-140 3-5 m 400-500 100-150 4-10 Al-ti-12& o.m 800-700 100-150 1-5 m 800-700 150-200 10-14

~-10_ 0.2-0.3_ 800-700 100-150 1-3

Al-H-121a. 0.5""" 1-2.51.Cu 700-800 150-200 1-3

Al-St>12B+_ Cu. Ni. etc 700-900 100-150 0-3

450-550 450-550

100-550 5-10 120-200 a-12

800-700 150-220 4-8 800-700 200-250 10-15

1700-800 150-250 2-8

laOO-l00 200-300 1-5

1800-100

1800-130

laOO-l00 100-200 . _ 0-2 _ Jl000-180 .) ArtifIrieIIy .. nom ~ DOt mdfted m = mtlfted. "1). UTS. in 1IN/m2

250-320 1-6

250-4C)O 4-0

200-300 0-3

.. o

I.4 - Machanical propartias and Microstructure of Aluminum-Silicon Alloys.

11

In general, the mechanical properties of aluminum silicon

casting alloys are strongly dependent uponJ•16 .:

1. The casting process used.

2. Dendrite arm spacing (OAS).

3. Grains of the primary phase.

4. Eutectic microstructure.

5. Degassing

6. Heat treatment.

These factors, in reality, are interconnected and often

interdependent in a complex manner. The interdependency

varies from one alloy system to another and from one casting

process to another. The casting process used dictates the

mechanical properties. The rapid cooling rate produced by

pressure die casting yields a fine eutectic microstructure,

smali dendrite arm spacing and smail grain size of the

primary phase, resulting in improved tensile strength and

elongationl5·17.18 (Table 1.4). Another researcher has reported

that reducing the OAS from 100 J.Lm to 25 J.Lm results in

increasing significantly the elongation from 2% to 13%, and

the ultimate tensile strength from 250MPa to 280MPa. From this

point, we can say that DAS obviously affects ductility of the

aluminum matrix.

other techniques are employed to achieve fine structures

or small grains. Reduction in grain size is brought about by

•

12

the efficient heterogeneous nucleation of the alpha aluminum

phase. This can be achieved through crystal multiplication

using mechanical or fluid flow forces to detach dendrite

arms l6• Practically, however, chemical additives are preferred

to refine the grain size by providing the necessary nuclei.

Titanium from 0.02 to 0.15%, or a titanium-boron constituent

(0.01-0.03%Ti and O.Ol%Br) is most widely used to reduce the

aluminum grain size, particularly for alloys with low silicon

concentrations l9• The TiA13 and (AI,Ti) B2 phases provide good

nuclei for a-aluminum lead~ng to finer grain sizes. A smaller,

equiaxed grain size is more preferable, as it improves the

resistance to hot tearing, decreases porosity and increases

mass feeding3,20.

Degassing is performed in the aluminum foundry to remove

the hydrogen present in the melt. By nature, hydrogen is very

soluble in aluminum, especially at high temperatures21 • l'he

presence of hydrogen, especially more than 0.30 mL/100 gr Al,

is associated with interdendritic porosity and secondary

(micron size) porosity. Hydrogen precipitation may alter the

form, as well as the distribution, of shrinkage porosity in

poorly fed sections of the casting3 • It is observed that, in

Al-l2%Si alloys, doublinq hydrogen from 0.2 ml H2/lOO gr Al up

to 0.4 ml H2/lOO gr Al results in a porosity increase from 0.4%

to 1.6%. Hence, hydrogen adversely affects the mec:'1anical

properties. The dissolved hydrogen in the melt, however, can

r

f

r [

)

13

be removed by fluxing with dry, inert and chemically pure gas

such as ni trogen, argon, chlor ine and Freon22•

As explained previously , in certain casting processcs

low cooling rates are encountered, such as in sand castings or

permanent mold casting. As a result, cast products with coarsc

microstructures are obtained. In other cases, for aluminuIn

casting alloys with silicon greater than 8%, the dendrite ùrIn

spacing (DAS) and grain size of the aluminum phase are not too

important, because the properties are mainly controlled by the

amount and the structure of the silicon phase7• Hence, in those

alloys, DAS and grain size play a minor role in the mechanical

properties. The microstructure of aluminum silicon eutectic

alloys is dominated by the eutectic mixture. The morphology i5

characterized by acicular flakes of the eutectic silicon

phase. The sharp edges of these particles act as a stress

raiser so that a crack will initiate and propagate easily. It

is necessary, hence, to refine the silicon phase morphology in

order to develop optimum mechanical properties. This can be

achieved by introducing small amounts of foreign atoms to the

alloy during the melt treatment called 'modification,l.

lModif ication, according to the ASM Handbook, is treatment of molten hypoeutectic, eutectic or hypereutectic aluminum­silicon alloys to improve the mechanical properties of the solid alloy by refinement of the size, shape of the silicon phase. Involves additions of small Wt% of sodium, strontium and phosphorus.

14

I.S - Modification

Since the present work is indirectly related ta

strontium modification, it is necessary to discuss the

modification phenomenon in a more specifie manner.

Modific:ation transforms the acicular eutectic silicon

phase to a fine lamellar or fibrous one, resuiting in improved

mechanical properties, especially elongation and impact

strengt.h2J.2~ 25. Figure 1.2 shows three different micro­

structures of the eutectic phase as a function of chemical

modification and casting process (solidification rate) related

to mechanical properties obtained for heat treated Al-si

hypoeutectic alloys.

Modification of the eutectic silicon phase was

discovered about 70 years aga. In 1921, Aladar Pacz found

that Al-si alloy containing between 5% and 15% silicon could

be modified by sodium fluoride to yield an improvement in

ductility and machinability~. The modified alloy since then

has been established commercially.

Modification can be accomplished by using a number of

elements such as those in groups lA and lIA in the Periodic

Table, as weIl as the rare earth elements27•28 • Among aIl of

these, however, only strontium and sodium impart a significant

modifying effect at low concentrations. The concentrations are

dependent on the silicon content of the alloy being treated,

the cooling rates, melting, as weIl as casting practices.

15

A. Acicular B. Lamellar c. Fibrous

~~~ ,ec ~ .. w~\{~ . 18 , •. " ':!11{.' ~I'".

',,"" ~ '" . :~, ....... ~\\',. ~ \. ~ " ~ .. -" ~,' '::

~l~-,jit~ ..... ft. • ~.. >

>J.~"" • : :~ 'If .': .. ,.~~_ ~ ~ ~ '. "'.1'<

.. " = .... 'Iol.' '\,_

1 ~ - . ~ ( '.' .~~ '.; ~ ~~, y'. -~.

" .(1').,~~:'~' .. -:,'. r~ .. .. ~~I ~'~ ~ t', ,IJ,/'f..... • .. l "

~,.. " . ,..,' U';1 ~. ,l . 'r .~~.~~ 'J ~~ t.. " ~ ~ ... if, ~ ~~.'1~. "", -: .: '..... ., ~,,~Vt~s~." 'i1J~'; ~ ~1 ,

• t t .,~ '.,~ r. ,;. t: '. :. l'~' <I;~ " · , ., ...... f'. ', •. '-'. ~ jl"~#Ô.ll'~l\1s. '. i!J; \.' .. ,.~. • l ~~ ~.. 'J j' ... ~ ..... ~ j.l" ~ 1'" ; ~,.~ ~

.-~. ". 'i " '" .:\~\~~>.~ .. ~ ,t!>'A"" ( ,1' " .-':. ,-~~ ~, \ t .... t" ..

;~\\t; "11-, ' I. .. ~~(" , ~ ~ "~~ : }.\.V'~'JT'" '" "~" ~. .... ' . 'i' ~\."" l 't.(:.( ':tIti

,,- ~ of.. &'" ' \; • ~~., .... ,''4 .. '~''''., i"'i- (, r .. ' (,tr- f "I .. { . , d;~~i, ~ . .,.. .. 4'*.. •• "' ~ ~ ... "_ , . ,"r',J,1. "'~~ Il'. •• ~ .. :hJ /~, "( .'~). ,', '\ . \" .

" " .. 1 .:-.JI'j :}.'.],.J t ".h • • 1/: },~ ...... \" ~ , :::" "'Î"-. ,,' ~

Al-7Si-o.3Mq Al-7Si-O.6Mg

Permanent mold

*) U'l'S E

MPa %

A 290 12

B 290 17

C 290 17

*) structure A = Acicular B = Lamellar C = Fibrous

SaneS Permanent mold

U'l'S E U'l'S MPa % MPa

275 2.5 330

280 .. 330

280 , 330

(unmodified alloy) (partially modifiad alloy) (well ModifieeS alloy)

E

%

10

13

13

Sand

UTS E

MPa %

290 2

295 2

285 2

Figure 1.2 Relationship between mechanical properties and

the eutectic microstructure which is a function of modification and casting process in heat treated Aluminum-Silicon alloys (24).

.

16

The generally accepted concentration range is 0.01% to 0.025%

for strontium8•

It is useful to discuss the different characteristics of

the major modifying agents that are commercially available.

Sodium modifies the silicon phase from a sharp-acicular

structure to a fine-fibrous one. It has two effects

eliminating aluminum phosphide (AlP) nuclei and rendering the

eutectic structure fibrous 29• The major shortcoming of this

element is that its action is only short-lived16 • This is due

to the high volatility of sodium at the holding temperature of

the molten metal. A further consequence of using this element

is that one cannot remelt and recast the modified material

since this causes a very significant sodium loss.

strontium has a similar effect to that of sodium; it

causes the eutectic structure to become fibrous30,31.32.

Quantitatively, the concentration of strontium required is

greater than sodium to yield the same effect. However, the

main benefit of strontium is that its rate of 10ss is

distinctly less than that ,';1" sodium, because Sr is less

reactive with the environment2. In many foundry practices,

strontium is the most common modifier used for aluminum­

silicon hypoeutectic alloy modification. commercially, it is

available as the pure element and as several Al-Sr master

alloys, such as; AI-10%Sr, Al-14%Si-10%Sr or Al-90%Sr with

different dissolution temperature characteristics ranging from

17

670°C to 750°C. Its stability in the melt enables the modified

metal to be remelted and recast without appreciable losses'.

The major drawback of strontium is that strontium-rnodif ied

alloys are more susceptible to porosity formation, especially

for sand castings with thick sections which solidify slowly'.

Based on a general foundry observation, it has been reported

that modified castings contain significantly more porosity

than unmodified alloys33.34.3.5.

Antimony is capable of changing the eutectic micro­

structure, but in a different way from that of sodium and

strontium it yields a lamellar structure instead of a

f ibrous oneS•36 • The degree of ref inement, moreover, is a

function of the solidification rate. With higher cooling

rates, antimony produces a fine lamellar silicon phase.

Antimony treated alloys can exhibit identical mechanical

properties to the same alloys modified with strontium or

sodium. The principal advantage of antimony is that its effect

is absolutely permanent. It is affected neither by holding nor

remelting3 • It is reported that antimony interacts with sodiurn

and strontium, which may be detrimental to the alloy ; thus,

it is essential not to mix the different modifyinq agents in

the foundry36. The major disadvantage of using antimony is

that it is toxic. Therefore, for hygienic and toxicological

reasons, antimony has not been widely used5•

, , •

18

In the 1960' s, i t was found that phosphorus is able to

refine the primary silicon in hypereutectic alloys. This

element facilitates nucleation of the silicon crystals by

enhancing the amount of the AIP compound which is naturally

present in Al-Si alloys. Phosphorus iR added to the alloy

in the range of 0.01%-0.03%. Phosphorus and sodium tend to

interact when both are present in the melt and render one­

another ineffecti ves. For hypereutectic alloys, sodium should

be kept at as low a concentration as possible to obtain refi-

nement of the primary silicon by phosphorus. Similarly, for

hypoeutectic alloys, phosphorus should be maintained as low as

possible to obtain good sodium modifications.

The modification is finally affected by various metal­

lurgical factors such as ; type of modifier, impurity pre-

sent in the alloy, cooling rate, casting method, method of

addition, addition temperature, degassing sequence, holding

time, type of degassing agent, superheat, pouring temperature,

fluxing, remelting cycles, and the silicon content37.38.39.4o.41.,n.

practically, impurity modification is often used in combi-

nation with quenching (in chilI casting). This technique

yields the finest microstructures.

The mechanism of modification has been the sUbject of

extensive research over the years. Several theories have been

proposed to answer these questions; however, it was not until

the use of transmission electron microscopy and x-ray

19

diffraction that the modification mechanism was clearly

understood. From various theories proposed, they can be

categorized into two groups43, namely, theories based on

nucleation mechanism and theories based on growth mechanisms.

The modification of the aluminum silicon alloy was

initially associated with restricted nucleation by the foreign

atoms. Crosley and Mondolfo44 have found that the nucleation of

silicon is controlled by aluminum phosphide (AIP) inter­

metallic compound. Since AlP has a crystal structure similar

ta that of silicon, i tacts as an effective nucleant. In

modified alloys (e.g.sodium), sodium neutralizes the

phosphorus, so that eliminating the nucleation by AIP and

increasing the undercooling for nucleation of silicon by

aluminum results in fibrous structures. It was assumed that

silicon particles are discontinuous~. This the ory has severai

weak points, since the silicon phase is continuous in the

modified, as weIl as in the unmodified, alloy, suggesting that

there is no repeated nucleation taking place after each

eutectic cell is formed2.4~. Moreover, it was found that

nucleation of the silicon from the bulk liquid does not oceur

in the modified alloy due to the presence of the nucleation

barrier46 • It is reasoned that, if the change in nucleation

kinetics was the only mechanism involved in the modification

process, only refinement of the silicon crystal would be

observed, but not the alteration in the shape of the

20

crysta147 • Therefore, the modifying agent must also influence

the growth kinetics of nucleated silicon particles during the

eutectic phase transformation.

In an unmodified alloy, it has been observed that the

silicon phase grows in advance of the aluminum phase with a

non-planar, non-isothermal solid-liquid interface. This leads

to the unbranched, flat-plate morphology·M. The single

crystals grown from the melt are bounded by definite crystal­

lographic planes, namely {111} planes. This plane is able te

grow only in a very specific direction, <112> direction. The

very important feature of silicon crystallization is that

twins are easily formed across (111) planes, resulting in a

self-perpetuating groove of 141 degrees at the solid-liquid

interface. This typical growth condition is called the twin­

plane re-entrant edge (TPRE) mechanismso,sl,s3.

The effect of modification would seem to be the

poisoning of the re-entrant {ll1} twin groovesSI ,S2.S3. Kobayashi,

et.alS4 have proposed that the modifying agent reduces the

anisotropic growth habit of silicon by adsorbing on specifie

crystallographic planes. In this way the modifier operates to

el!4ilinate the silicon lead distance and to permit semi­

isotropie fibrous growth similar to that observed in chilI

modification. Moreover, L.M Hogan et.alsl have observed that

modif iers produce an enormous increase of twinning in the

eutectic silicon. This increase has been associated with the

..

'."

21

reduction of the twin boundary energy, thereby increasing the

growth flexibility of the faceted phase. By a careful study

using transmission electron microscopy, Lu & He lawe 11''

finally found that only impurity atoms with a suitable size

are capable of modify ing the si licon phase. l t ';Jas establ ishecl

that geometrically the ideal radius ratio r,,,,, .. hli<, to r"h"",' is

about 1.646. The impurities are adsorbed on the silicon solid

liquid interface. Their atomic positions promote twinning by

displacing a {Ill} mono layer growth step to an alternative

stacking sequence. This concept is then called "impurity

induced twinning", and i t has been widely accepted as the best

approach to explain the modification phenomenon.

I.6 - The Role of Thermal Analysis in Modification Assessment

Thermal analysis has been used to study the solidifi­

cation of metal since the beginning of this century. Th i5

method was mainly used to establish the equilibrium phase

diagrams for metals and alloys. Historically, the use of

thermal analysis to study the evolution of microstructure in

aluminum casting dates from the publication by Cibu la et. a l VI.

Their study on Al-Cu alloys established that a cooling curve

can be used to determine whether or not the alloy i5 either

grain refined or rnodified. Grain refinement can be assessed

by looking at the change of undercooling in primary phase

· i

.

.--- --------------------------------------

22

(Al) crystallization. Furthermore, by studying the

solidification of Al-Si alloys, Crossley and Mondolfo57 showed

that modification affects both the nucleation and the eutectic

temperature. As a result of the extensive development of

microprocessors, it is now possible to automate the recording

and analysis of the cooling curve. Thermal analysis as a non-

destructive test for assessing modification has been largely

used in combination with conventional meta110graphic examina-

tion. Current1y, it is reported that thermal analysis equip­

ment is used in hundreds of aluminum foundries around the

world to control the microstructure prior to casting58 •

The effect of modificationS9,60,61 on the co01ing curve is

reported to cause first of aIl, a depression of the eutectic

plateau, secondly, an increase in the undercooling required to

start the eutectic freezing 1 and f inally, a lengthening of the

undercooling. Both the nucleation and growth temperature of

the eutectic silicon are lowered when the alloy is properly

modified (see Figure 1. 3). The nucleation point is 10wered

since the mOdifying agent, strontium for instance, neutralizes

the AlP compound at sites where the silicon crystals start to

nucleate. The eutectic growth temperature ia also affected

because the strontium plays a significant role in the eutectic

silicon growth mechanism. From the three features listed

above, the eutectic depression (5°- 10°C) is the easiest way

to assess whether or not the melt has been proper1y modifiedi

23

therefore, the depression of the eutectic temperature is the

most commonly used method to evaluate modification5l!.o~.

-U -

T e

Figure 1.3

"':"'T ·1 .. ··· 1

'te = Eutectic temperature

-~ T = Undercooling

-Te= Eutectlc depression

Unmodified

Modlfied

., .. t. -~ T e

------__ .~ Tlme

The alteration of the cooling curve due to modification [59].

When dealing with the cooling curve, we sometimes find it

ls difficult to locate exactly the important inflection points

on the curve. To overcome this problem, a derivative cooling

curve (dT/dt vs Time) can be employed as first described by

Backerud in the 1970s63• The main advantage of the derivative

cooling curve is its ability to magnify inflection points

which are hardly visible in the ordinary cooling curve.

Figure 1.4 indicates several important points both on the

cooling curve and on its derivative. The arrow locates the

silicon nucleation starting point. Many researchers are

f ,

.. 1 ...

r

24

reported to have misinterprated the silicon nucleation point

by referring to the minimum undercooling temperature58.62

•

650 6 .. 0 630 620 610 - SOO (J

"" • 590 ... , 580 -0

i 570 r 560 .- 550

5 .. 0 530 520 510 500

0

Figure 1. 4

0.5

O." 0.3

0.2

0.1 Q, .... " 0 Q. ---{).1 n

" Il

-{).2 • n ...... -().3

-{) .. -{).5

-{).S

-{).7 50 100 150 200 250 300 350 400

r.,. (YCOnd)

Cooling curve and its corresponding derivative curve for Al-7Si-1Mg synthetic alloy [63].

I.7 - Quantitative ADalylis and Silicon Partiel. Charaeterization

Although quantitative metallography has been used for a

long time in the characterization of particles, its applica­

tion to the eutectic silicon morphology analysis in Al-si

foundry alloys dates from perhaps a decade ago. The reasons

:)re probably due to the broad application of modification only

sinee the past ten years, and also due to the recent develop-

ments in electronic data proeessors, so that one can eollect

and process hundreds of data rapidly. Tuttle and McLellanM.

2S

have characterized the silicon particles in A357 alloy as a

function of solidification rate and heat treatment. They also

have made an attempt ta corre1ate the silicon features to the

mechanical properties, especially ductility. It was observed

that the silicon partiel. taret diamatar2 decreased linearly

as solidification rate increased. There is an increase of the

partiel. count at higher solidification rates. This is rnerely

due ta chilI modification. Moreover, increasing the solution

treatment time and temperature as well as the freezing rate

result in a decrease of the aspect ratio leading to particle

sphericity. It was found that the ductility is inversely

related to the aspect ratio whereas tensile strength is

directly related. Similar results are also reported in

reference 65. More recently, S.Shivkumar, L.Wang & D.Apelian

have observedM, that in a modified 356 and 319 casting alloy

wi th a certain cooling rate, increasing strontium beyond

0.015% causes the average diameter and aspect ratio of the

silicon particle to increase (see Figure 1. 5). This coarsening

of the silicon particle, in addition to the porosity factor,

perhaps explain the reduction of the mechanical properties at

strontium higher than 0.015% which was observed by Closset and

Gruz leski 23.

2According to M. Tuttle et. al [64], 1Ilaxi1llUlll teret diamater is defined as the maximum linear dimension, units are in Mm, count is the average number of particles per field, and aspect ratio is the ratio of maximum to the minimum dimensions

~ "

(

{

26

4 30

o Alpect Ratio 29 • Ollmet.r

_ 3 28 j

! ... 2 7 ~ s 1 al i! 2

2.5

2.4 0.0 001 0.02 0.03 004

% Strontium

Figure 1.5 The variation of silicon partie le diameter and aspect ratio with strontium concentration [66] .

The parameters used for silicon partiele characteri­

zation carried out in the current work will be further

discussed in Chapter II.

I.8 - The Aims of The Present Work

As previously described, in heat treatabl.2 aluminum-

silicon alloys, magnesium is introduced in small percentages

This ranges trom 0.35 wt' (A35ô) to about 0.55 wt% (A357).

The reason for magnesium addition ls to strengthen the

aluminun matrix by precipltatinq a hard-brittle MqzSi phase

during the heat treatment.

The exact effect of magnesium on the eutectic

microstructure, however, has never been elarified, and

confusion exists in the literature. Gruzleski and Closset3

found that, in commercial alloys with hiqher magnesium

27

contents (A357), the eutectic phase is somewhat finer than in

alloys with lower magnesium levels (A356). On the other hand,

Bercovici~ has reported that magnesium coarsens the eutectic

microstructure for the same alloys.

Further addition of magnesium leads ta the formation of

a ternary eutectic phase. Numerous papers have discussed the

Al-si-Mg phase diagram frem the thermochemistry and thermo­

dynamics points of view1.61

• Sa far, however, little attention

has been paid to the role of the magnesium content on the

eutectic microstructure. If certain levels of magnesium are

indeed able to refine the microstructure, it couid possibly be

used as a modifier. By contrast, if magnesiurn coarsens the

microstructure, which is undesirable due to negative effects

on the mechanical properties, it should not be used in

excessive amounts. It is necessary, therefore, to f ind thF.!

real effect of magnesium on the eutectic microstructure and to

study its role, if any, in the silicon nucleation and growth

mechanisms. Hence, the purposes of the present rese3rch are ;

1. To discover the role of magnesium on the silicon

microstructure of unmodified and modified Al-Si casting

alloys.

2. To discover the possible role of magnesiu~ on the

solidification of silicon in Al-si casting alloys.

3. To study possible interactive effects between magnesium

and strontium in Al-si-Mg ternary alloys.

28

Chapter II

EXPERIMENTAL PROCEDURE

2.1 - General

To ascertain the role of magnesiurn on the eutectic

microstructure, certain synthetic Al-Si aIIoys had to be

produced. These aIIoys can be categorized as hypoeutectic

nlloy, that is, Al-7%Si, and eutectic alloys, Al-ll%Si, and

Al-13%Si. The influence of magnesium on the silicon morphology

was aiso investigated in the presence of strontium, a well-

known foundry modifier. The arnount of magnesiurn used in this

study was va~ied from 0.05% to 1%, whereas strontium was

kept around 0.020% - 0.025%. In addition, Mg was aiso added

to certain widely used commercial alloys such as 413.2, 356.2

and 357.2 foundry alloys. The total number of samples studied

was about 42. The silicon morphology, strongly affected by the

modification and casting parameters, was studied using thermal

analysis and also quantified with an image analyzer coupled

to a scanning electron microscope. The entire exper imental

procedure can be seen schematically in Figure 2.1.

2.2 - Alloy Preparation & Alloy Makinq

2.2.1 Materials

In order to eliminate contaminants, synthetiG alloys

were prepared trom high purity constituents.

,r,

Aluminum

COMPUTER

THERMAL ANALYSIS SYSTEM

COOLING

CURVE

ANALVSIS

MATERIAL PREPARATION

AS CAST MATERIAL

INDUCTI~ FURNACE 1

SAMPLE PREPARATION

'---___ Ir 1

OPTICAL

METALLOGRAPHY

IMAGE ANALYSIS

Figure 2.1 Experimental Procedure

29

Aluminum, the main ingredient used in this experiment, was

supplied by Belmont Metals Co. This is 99.99% pure metai in

pellet forme

30

sil, ieon

The semiconduetor-grade silicon (99.999%> in lump form was

purehased from JOhnson-Mathey Inc.

Magnesium

This metal, about 99.50' qrade, was supplied by Ameriean

Chemical Ltd in flake forme The composition of the metal is

as follows :

Mq Cu Fe Pb Mn Ni Sn wt% Wt' wt' wt' wt% Wt% wt%

95.50 0.02 0.035 0.01 0.15 0.001 0.01

Strontium

strontium was obtained from the commercial master alloy, Al-

10%5r. The composition of this alloy allows the maximum Sr

dissolution in molten aluminum silicon alloy at 750°C.

The remùininq material, 413.2, 356.2 and 357.2 alloys,

were purchased as commercial inqots. Their compositions are

qiven in the Appendix.

2.2.2 .eltiD9 and ca.tiD9

Three pieces of equipment were used in alloy preparation:

1. Induction Furnace where the alloy was fabricated.

2. Ceramic crucible, as a mold, into which the molten metal

was poured.

3. Thermocouples connected to a microcomputer.

.'

31

Appropriate amounts of aluminum and silicon were first

melted together in a 6 cm diameter ceramic crucible placed

inside a copper coil in the induction furnace. The furnacc,

capable of producing power as high as 100KW, is made by Tacca

Ltd. The Al-si metal was heated electromagnetically and

melted for just 10 minutes. The molten metal temperature was

maintained in the range of 740°-750°C. A measured amount of

magnesium wrapped in aluminum foil was then introduced, using

a stainless steel plunger. To minimize magnesium burn-off in

contact with air, an inert gas, argon, was blown continuosly

across the surface of the molten metai. The system was held at

constant temperature for about 10 minutes to dissolve the

magnesium completely, and the alloy at 750°C was then cast

into a smaller crucible • The liquid alloy was then allowed to

sOlidify, and during that period, a chromel-alumel K-type

thermocouple, of O.Ol°C sensitivity, was used to measure the

ternperature. Thermocouples were placed in the center and the

edge of the mold to record the temperature gradient between

the two points while the metal was freezing. To avoid direct

contact with the hot metal which could damage the junction,

the thermocouples were sheathed in 0.044 inch diameter

stainless steel.

In the case of magnesium treatment in the presence of

strontium, prior to pouring, a small amount of strontium

(0.02%) was added to the Al-Si-Mg molten alloy at 740°-750°C.

i ,

32

The system was then held for about 5 minutes to dissolve the

strontium completely. To improve the homogeneity, stiring was

performed using a steel wire coated with ladle wash. Figure

2.2 shows the casting set-up, including the thermal analysis

apparatus, and Figure 2.3 shows the molten aluminum-silicon

alloy.

2.3 - Thermal Analysis

Thermal analysis was performed to study the cool1ng

curve which reflects the nucleation and growth process of the

eutectic phase. Once the thermocouple registers a temperature

change at any given time, it transfers the data electrically

to the microcomputer. The microcomputer operates with thermal

analysis software manufactured by Omega Engineering Co. The

system was set for 0.025 mV resolution with low noise, and the

time per channel was 20 milliseconds. The delay time between

inputs was fixed to be 2 seconds. The data was collected and

stored on diskettes to be further processed using LOTUS and te

be pletted , as cooling curves, using a Hewlett Packard

pIotter. A cooling rate of about 2° C/second, was generally

obtained in the experiments using a ceramic crucible as a

mold.

Derivative thermal analysis curves (dT/dt) were aiso used

to locate the eutectic growth temperature and te locate the

possible reactions leading to intermetallic phases in the

structure (ternary alloy).

Figure 2.2 casting and Thermal Analysis System

Figure 2.3 Molten metal of Al-Si alloy where rnagnesiurn and strontium were introduced.

33

p

i ,

34

2.4 - spactrochemical Analyais

Chemical analysis was carried out to determine the actual

composition of both commercial and synthetic alloys. The

analysis was done using an optical emission spectrometer

manufactured by Baird Atomic, Spectro 1000, model No. DV2

interfaced with a microprocessor. The analysis results are

listed in appendices to this thesis.

2.5 - Metalloqraphie Examination

As-cast samples, about 7 cm high and 5 cm in diameter,

were cut to appropiate dimensions as shown in figure 2.4 and

mounted in Bakelite resin. The surfaces were successively

ground with 120, 280, 400 and 600 grit silicon carbide paper.

After ul trasonic cleaning, the samples were then polished

using a 5 ~m and 0.3 ~m alumina powder aqueous suspension on

a cloth-covered wheel and finished using colloidal silica

solution. A selective etching with 5% NaOH in water was done

to improve the contra st between the aluminum matrix and the

silicon phase, particularly for scanning electron microscope

examination.

2.5.1 optical Microscopy

Polished samples were examined under a Neophot optical

microscope. They were then selec'ci vely photographed using

polaroid 52 film. The micrographs obtained are presented in

Chapter III.

-.

, 1 ~

-'1 2.5 cm 1

f--7cm

1 1

~ ......

:~ 1

As cast sample

Chemical Analysis

Metallography

Examination

Fig 2.4 Sample Preparation

36

2.5.2 scanning Electron Microscopy

Selected samples of hypoeutectic Al-7%Si alloy and

hypereutectic Al-13%Si alloy were etched and observed under

a JEOL-JSM 840 scanning electron microscope shown in Figure

2.5. This versatile instrument is capable of producing an

image with high resolution. Images displayed on the CRT may

come from either secondary electrons or backscattered

electrons captured by the detector. The accelerating voltage

was set for between 2~-35 kV, and the working distance (WD)

and probe current were fixed at 39 mm and 3 x 10-10 Amp

respectively. Several pictures were taken at various

magnifications ranging from 300 X to 600 X. In the case of

intermetallic compound observation, the SEM was set up in the

EDS program and the image obtained on the CRT originated from

X-ray emissions captured by the LN2 X-ray detector. The

chemical stoichiometry of intermetallic compounds was

determined using WDX with microprobe analysis.

2.6 - Quantification Using Image Analysis

Silicon particle characteristics were defined using image

analysis. In this thesis, the particle features as defined by

the Northern.Co are adopted68• Figure 2.6 illustrates how the

particles are positioned and measured in the X-y coordjnates.

This figure also describes the various size and shape

parameters used.

\

i ~

, ~ ~

~ 1 ( )

~

! ~

~

/

37

Figure 2.5 Scanning electron microscope coupled with an image analysis system.

f

38

vu Cent.r of partiel.

t -.----------

Qi ... .f -. __ .-. ----., >-

1 1

+-1 1 1 1 1 1 t 1 1 1 1

1

---_.!._----- 1 1 • 1 1 1 1 • t 1 1 1 -0 1 X '''- X·Feret ......,

PARAMETER DEFINITION

The distance around the edqe of the 1 perimeter partiel •• If the partiel. has a hole in

it, P also inclueSes the distance around the 8dqe of the hol.

Shape Factor perimetero/(4'1' Area)

x-reret Projection of the partiele on the x-axis

Y-reret projection of the partiel. on the Y-axis

Aspect Ratio Lenqth 1 lfidth

Area The aua of all pixel composinq th. image of the partiel.

Figure 2.6 Partiele eharacterization. The partieles are first positioned in the X-Y coordinates, the x and y ferets are then projected onto the axes, and finally, a number of parameter are defined and calculated [68].

"

..

"

.'

39

Image analysis was conducted with a pair of image

analyzers, TN5700 & 5400, produced by Tracor Northern Co. They

were interfaced with the SEM through a program file, cailed

IPA57. The system was able to extra ct the video image on the

SEM to he further proccessed. In principIe, the main stages of

image analysis are (see figure 2.7) ;

1. Image acquisition

2. Binary creation

3. Object detection

4. Measurement

A brief look at each will be taken. The video image

obtained from the SEM is first transformed into an electronic

signal suitable for processing and storage. This digitized

signal is known as the gray level image. The gray levei image

actually consists of thousands of tiny dots known as pixels or

picture elements. Pixels are used to define image resolution,

that is, the more pixels required, the higher the resolution.

Image acquisition is a primary concern of image analysis;

therefore, it is important to have images with equally good

contra st and resolution. Once this is obtained, the image is

not succeptible to degradation, and it can be kept in storage.

The next step is binary creation. This operation separa~es the

region of interest (silicon particles, for exarnple) from

background (aluminum matrix) by binary black & white image.

The binary image, derived from the grey level image in mernory,

(

(

(

40

creates pixels in ON (full intensity) or OFF (black). Binary

images can be analyzed as is, but to obtain more accurate

da ta, images are further improved through fil ter ing (eg ;

dilatation, erosion, edge). This operation can enhance the

image definition by separating contiguous particles or fused

stringered features~.

The third step, object detection, is to define each

region of interest (ROI) within one particular image. This

step allows data reduction by guarding the image and

establishing the number and dimensions of particles to be

measured. Object detection 1s capable of describing each

object more compactly than a l1st of coordinates of every

single pixel it encompasses. Once discrimination has been

completed, aIl necessary information can be used in subsequent

measurement operations.

The last step, the real interest in peforming image

analysis, is to carry out particle sizing or shape measurement

The technique is based on the number and distribution of

detected particles relative to the scan line and the total

number of particles. The particle size and shape analysis

allows one to characterize various features of the phases

present. In the current work, the parameters thought to be

most sensitive to the modification effect on the silicon phase

were defined. These are perimeter, shape factor and aspect

ratio. AlI those features can finally be proccessed and

•

, ,~

.cr ..

"""

41

presented as graphs and tables. All of the stages described

above are carried out using an interactive mode with IPA57 and

a more powerful data management program (TECHCALC).

IMAGE GREY IMAGE ACQUISITION PROCESSING

~-I GREY IMAGE

~ FILTERING

, 1 PRINTI

SET-UP PARAMETERS

~I MEASUREMENT

Figure 2.7 Image Analysis Block Diagram

,

.( 1

"

42

On peforming image analysis, errors may be introduced

from many sources, hence, several precautions have to be

taken. To be able to obtain a consistent resul t, i t is

necessary to set up the experimental parameters exactly the

same from one measurement ta another. This will increase

repraducibility. Image magnification also plays an important

role. As magnification increases, greater field-to-field

measurement variations are encountered. Thus, increasing

magnification improves resolution, but on the other hand, the

area measured may not be representative, particularly if the

number of fields is limited. Therefore, a compromise is needed

between the requirements for high resolution and those for

good statistical sampling .

.. Chapter III •

RESULTS , DISCUSSION

43

The results of the present work are presented in this

chapter as cooling curves, graphs, tables, and micrographs,

which are analyzed and interpreted in order to assess the

effect of magnesium on Al-si alloys. Thermal analysis was

carried out for Al-7%Si, AI-ll%Si and AI-13%Si synthetic

alloy, as weIl as for a 413.2 commercial alloy. The cooling

curves obtained for the Al-7%Si alloy are shawn in Figures

3.1 and 3.2. To deterrnine the eutectic temperature, derivative

cooling curves were used. These results are listed in Tables

3.1-3.2. The sarne method was also applied to the other alloys,

and the final results are presented in Tables 3.3 through 3.8.

The micrographs are arranged 50 that the evolution of the

silicon morphvlogy due to magnesium and strontium additions

can be clearly followed.

3.1 - Maqnesium' Strontium Effects in Hypoeutectic synthetic Alloy

Al-7%Si Alloy

The cooling curve for an Al-7%Si alloy can be seen

in Figure 3.1. When the melt is cooled at a cooling rate of 2°

C/s, it initially exhibits a primary phase solidification

(Al) as indicated by the first large inflection point. After

. '-

44

a period of time, this is followed by eutectic freezing

marked by the thermal arrest. with high purity Al-si alloy

solidified in a ~Brdmic crucible, the eutectic temperature is

found to be 57a.5°C, and the under~ooling is O.6°C (Figure

J.1, the first cooling curve). The temperature is stable for

a long period of time during eutectic growth, and then drops

rapidly on the completion of eutectic freezing.

By adding Magnesium, about 0.06%, the eutectic

temperature is shifted slightly (see Table J.1) to 57aoc from

578.5°C while the undercooling is changed to 0.5°C. The

micrographs in Figure J.4 indicate that there is no signifi-

cant effect on the silicon phase, although some change of

silicon shape can be observed in several areas of the

microstructure. By increasing the magnesium content from 0.36%

to 0.44%, the eutectic temperature is further depressed and

the undercooling, AT, completely disappears. The magnesium

causes the silicon morphology to change from a lamellar to a

fine acicular structure. Further additions of magnesium of up

to 1% decrease the eutectic temperature significantly by

17.6°C. This is accompanied by an increase in the undercooling

of about 1.5°C. The cooling curve of this high Mg alloy has an

inflection point at about 555°C which suggests the presence of

ternary compounds of Al-M92Si-Si. The final structure (Figure

3.4d) consists of a fine lamellar and possibly even a fibrous

structure, occuring in several small pockets. Bo~h the

t L ___ _

4S

morphology of the silicon phase and the depression of the

eutectic temperature (~TE)' as weIl as the increased under-

coo11ng suggest that high levels of Mg act to slightIy modify

the silicon phase in this particular Al-7%si synthetic alloy.

The addition of 0.02% strontium to the alloy depresses

the eutectic temperature by about 2.6°C (see Table 3.2). The

micrograph of this structure shows (Figure 3.5a) a dramatic

change in the silicon phase morphology from a larnellar to a

fine fibrous structure. The strontium addition emphasizes the

dendritic morphology. The dendrites, characterized by a dis-

tinct DAS, can now be clearly viewed. By introducing magnesium

to this alloy, the eutectic temperature was lowered, (Figure

3.2) and the effect of strontium modification decreased. This

effect i5 most pronounced at high magnesium concentrations

(Figure 3. 5d) .

, , j

AI-7%Si alloy Wlthout StrontIum modificatIon

650

640 ,-630

0.000/0 Mg

620 0.060/0 Mg

610 0.35% Mg

600 0..440/0 Mg -Q. 590r ~, UJ 1.01% Mg a: ::J

580 S UJ 570 Il. ~ w 560 1-

550

540

530

520

510

500 0 100 200 300 400

TIME (SECOND)

Figure 3.1 The cooling curves for unmodified Al-7%Si alloys. • Notice the shifting of the cooling curve shape as

en

magnesium levels increased.

--- - - --"--, _. , .~;:S~~!l!\~~~ ... ~~~-",,;i~-or-~'t ~.f;ic.~ ... ~-~~'''"'''-''''~''''''''''''''' "'"~ ..... ,--",~"'~"<-r

AL-7%Si With Strontium modification

650

640 0.00% Mg 630 ----- 0.06% Mg 620 0.41% Mg 610 J- ~ ----- 0.61% Mg 600 r ~

6 1.21% Mg

- 590 W a: 580 :::> ~

570 1 a: """ ,"x -.-.... ~-... _-UJ Il. ~ 560 w t-

550

540

530

520

510

500 0 100 200 300 400

TIME (SECOND) ~ ~

F1Dure 3.2 The cooling curves fer modified Al-7%51 al:oys.

,#

-~ ~

r

48

Table 3.1 The effect of magnesium addition on Al-7%Si synthetic alloy

Mg Sr TNl TN2 ATN TEl TE2 ATE (%wt) (%wt) (OC) (OC) ( OC) ( OC) (OC) (OC)

0.00 610.6 614.0 3.4 577.9 578.5 0.6

0.06 612.0 615.3 3.3 577.5 578.0 0.5

0.36 608.4 610.3 1.9 573.2 573.3 0.1

0.44 609.3 612.3 3.0 571.4 0.0

1. 01 607.6 611.2 3.6 559.4 560.9 1.5

Table 3.2 ."l- e effect of magnesium in the presence of strontium on Al-7%Si synthetie alloy

Mg Sr TN1 TN2 ATN TEl TE2 ATE (%wt) (%wt) (OC) (OC) ( OC) ( OC) ( OC) (OC)

0.00 0.030 612.2 613.8 1.6 573.4 575.9 2.5

0.06 0.030 612.5 614.0 1.5 572.5 574.0 1.5

0.41 0.028 610.0 613.3 2.0 567.9 568.5 0.6 0.61 0.029 610.0 611.7 1.7 563.4 564.5 1.1

1.21 0.025 607.5 610.5 3.0 556.6 558.0 1.4

Note Cooling rate observed in the experiment using eera mie mold is 2°C/sec. TN1, TN2 and ATN are the mini mum, growth and undercooling temperature, respecti vely for primary aluminum. TEl, TE2 and ATE are temp­erature for minimum, growth and undereooling temp­peratutre respectively, for eutectic phase. (See a schematic of a typical eooling curve of Al-si alloys in Figure 3.3)

r 49

Primary Aluminum Solidification

1 • T N --- ------ ------------------------2 ~TN

T N, ---- --------- --------------,-----

----------,---------------- -------~---~. TE2

,0. TE 1 TEt ---------r----------- --------------- ------- -- -

Eutecfic Solidification

-------.. ~ rime

Fig 3.3 A typical cooling curve for AI-Si AI/oys

r

a) 0.0 % Mg

Figure 3.4 The evolution of eutectic microstructure as a function of magnesium content in an unmodified Al-7%Si synthetic Alloy.

1 •

--------------50

b) 0.36% Mg

d) l%wt Mg

a) 0.03% Sr + 0% Mg

c) 0.029% Sr + 0.61% Mg

Figure 3.5 The evolution of eutectic microstructure as a function of magnesium ~ontent in strontium modified Al-7%Si synthetic Alloy.

51

cl O.03~ S o r + 1 _) 0.06% M _.! 1 :r 9 ,'4 r ,~_,.ro ~. ,-4 J..... .

• "t.' '~.,.-q~~. "'~~ / };~k , ~"'l~ • 1 il. . ..<J. -' • '\' \ IC.~

~' ~-..

" '.. ,'" , .. ,\; . .,...-... ~. _ ':; . ,.-;' ~aa • w" <~. . ,~lk,. -~ __ , J~..~ . ~ ',> \.: • .,..

+ 1. 21% Mg

(

3.2 - Maqnesium , strontium Eff.ct in an Eut.otio synth.tic Alloy

Al-ll%si Alloy

52

Essentially, there is no significant difference between

Al-ll%Si and Al-7%Si, as indicated by cooling curves and

microstructures. In the pure Al-lltSi alloy, the eutectic

growth temperature is 576.3°C while the undercooling is 1.SOC.

The silicon morphology is characterized by a sharp acicular

structure with several islands of primary phase. The addition

of a small amount of Mg slightly decreases the eutectic growth

temperature from 576.3 to 575.4 oC without affecting the

undercooling (see Table 3.3). Fu~ther additions of magnesium

of 0.34 % and 0.62% depress the eutectic temperature by 3.5

and 5. 2°C respectively. The micrographs clearly show the

alteration of the silicon phase from a lamellar to a coarse

acicular-faceted structure (see Figure 3.6). The islands of

primary dendrites gradually become bigger with increasing

magnesium, and the structure seems to conta in less eutectic.

The only difference between the Al-lltSi and the Al-7%Si alloy

is that, in the former case, the degree of refinement is

decreased, less than that observed in Al-7%Si alloy. In the

case of strontium modification, increas:ing the magnesium

content agaln coarsens the silicon phase (see Figure 3.7).

Here, again, it is observed that the magnesium masks the Sr

modification.

S3

Table 3.3 The effect of magneslum addition on Al-ll%Si synthetic alloy

Mg Sr TNl TN2 ~TN TEl TE2 ATE

(%wt) (%wt) ( OC) (OC) ( OC) ( OC) ( OC) ( oC)

0.00 577.6 582.0 4.4 574.8 576.3 1.5

0.08 580.0 582.8 2.8 573.9 575.4 1.5

0.34 579.9 583.2 3.3 572.0 573.2 1.2

0.62 579.0 582.0 3.0 570.3 571.1 0.8

1. 06 579.0 581.6 2.6 564.2 566.0 1.8

Table 3.4 The effect of magnesium in the presence

Mg (%wt)

0.00

0.08

0.33

0.60

1. 20

Note

of strontium on Al-l1%Si synthetic alloy

Sr TNl TN2 ~TN TEl TE2 ATE (%Wt) (OC) ( OC) ( OC) ( OC) ( OC) ( OC)

0.021 582.0 583.0 1.0 573.2 574.0 0.8

0.018 581.8 584.5 2.7 572.4 573.6 1.2

0.020 580.0 583.0 3.0 568.0 568.8 0.8

0.025 578.0 581. 7 3.7 564.7 566.8 2.1

0.021 580.5 583.4 2.9 562.0 564.0 2.0

Cooling rate observed in the experiment using cera

mie mold is 2°C/see. TN1, TN2 and ~TN are the

minimum, growth and undercooling temperature,

respectively for primary aluminum. TEl, TE2 and ATE

are temperature for minimum, growth and undercooling

temperature respeetively, for eutectic phase,

(consult Figure 3.3).

a) 0% Mg

c) 0.34% Mg

Figure 3.6 The evolution of eutectic microstructure as a function of magnesium content in an unmodified Al-l1%Si synthetic Alloy.

•

'.

54

100p.m

b) 0.08% Mg

l ,

100 p.m l. iii l "I~

d) 1. 06% Mg

'f 1

\

a) 0.021% Sr + 0% Mg

c) O.O?% Sr + 0.33% Mg

Figure 3.7 The evolution of eutectic microstructure as a function of magnesium content in strontium modified Al-11%Si synthetic Alloy.

b) 0.02% Sr + 0.33% Mg

d) 0.025% Sr + 0.6% Mg

55

Al-13%Si Alloy.

The influence of magnesium on the Al-13%Si synthetic

alloy is very much the same as that seen for the hypoeutectic

Al-ll%Si alloy. That is, at a high silicon content, the

magnesium changes the silicon phase as seen in Figure 3.10.

The depression of the eutectic temperature and the alteration

of the undercooling can be seen in Figures 3.8 and 3.9, as

weIl as in Table 3.5. Two interesting features are observed on

the cooling curve of the high magnesium containing alloy. The

growth of the aluminum primary phase now is more pronounced,

and the formation of the ternary phase at about 550°C is

visible. The series of micrographs (Figure 3.10) reveals that

as the magnesium content is increased, the eutectic morphology

evolves from a sharp-acicular-flake to a rough faceted

lamellar structure. At about 0.98% magnesium, the presence of

ternary compound~, AI-M92Si-Si, as evidenced by chinese-script

dark phases, is more pronounced. Indeed, at this high Mg

content, as found earlier in Al~7%Si and Al-11%Si, sorne degree

of refinment has occured in several area of the microstruc­

ture. Generally, however, the eutectic silicon is coarse.

The effect of magnesium addition on the strontium modifi­

cation is that the eutectic morphology becomes less and less

modified as the magnesium content is increased (Fig 3.11). The

final microstructure is a mixtuLe of fibrous, lamellar, and

acicular silicon. The image analysis detailed in a separate

section shows this diversity in a more quantitative manner.

r .

650

640

630

620

610

600

~ UJ 590 li: ~

~ 580

, UJ 570 r-~

~ UJ 560~ t-

\

550r 540~ 530~

1 1

520f-i ,

510 ~-1

500 L-

AI-13%Si Alloy Without StronIild Modification

o 100 200

TJUE (SECOND)

0.00% Mg

0.09% Mg

0.55% Mg

0.81% Mg

0.98% Mg

------- --- --'-----300 400

Figure 3.6 The cooling curves for unmodlf~ed Ai-:7~S_ 7~!_:~ Notice the shlfting cf tne COCll ..... ] ,_urVE:: :"'I<:j_=_ ,-j:

magneslum level lncreaseà.

: .. ,. 'J

,-, AI-13%Si Alloy Wt1h Strontilm Modification

6501 l ---------- _. - -,

640 ~

~ 0.00% Mg

1

630~ ---- 0.08% Mg 1

620~ 0.57% Mg 610 ~ Il -~- - 0.70% Mg ID)~ 1\ 1.40% Mg

6' 590 -w a:. 580 :::l

i 570 w 0... ::::t

560 w t-

550

540

530

520

510

500 0 100 200 300 400

TlME (SECOND)

'" CD Figure 3.9 The cooling curves for modified Al-13%Si alloys.

S9

Table 3.5 The effect of magnesium addition on Al-13%Si

synthetic alloy

Mg Sr TN1 TN2 ~TN TEl TE2 ~TE

(%Wt) (%wt) ( OC) ( OC) ( ° C) ( OC) ( oC) ( OC)

0.00

0.09

0.55

0.81

0.91

Table 3.6

Mg (%wt)

0.00

0.08

0.57

0.70

1.40

Note

564.5 575.5 11. 0

565.4 575.2 9.8

566.0 573.5 7.5

563.2 571.6 8.4

563.3 566.7 3.3 56~.6 567.6 2.0

The effect of magnesium in the presence

of strontium on Al-13%Si synthetic alloy

Sr TN1 TN2 ~TN TEl TE2 ~TE

(%wt) ( OC) ( OC) ( OC) ( OC) (OC) ( 0 C)

0.024 566.1 574.3 8.2

0.030 565.5 573.1 7.6

0.028 569.8 570.0 0.2

0.030 567.6 568.6 1.0

0.034 563.2 565.0 1.8

Cooling rate observed in the experiment using

ceramic mold is 2°C/sec. TN1. TN2 and ûTN are

the minimum, growth and undercooling temperature

respectively for primary silicon phase. TEl, TE2

and ~TE are the minimum, growth, and undercooling

temperature respectively, for eutectic phase.

L

r

al 0% Mg

c) 0.55% Mg

Figure 3.10 The evolution of eutectic microstructure as a function of magnesium content in an unmodified Al-13%Si synthetic Alloy.

60

d} 0.98% Mg

..

a) 0.02% Sr + 0% Mg

c) 0.03% Sr + 0.57% Mg

Figure J.11 The evolution of eutectic microstructure as a function of magnesium content in strontium modified AI-IJ%si synthetic Alloy.

,

, , .

L.

61

d) 0.034% Sr + 0.70% Mg

3.3 - Magnesium' strontium Effects in 413.2 commercial Alloy (10.55%8i)

62

The data obtained from the cooling curves of the 413.2

commercial alloy are listed in Tables 3.7 and 3.8. Due to the

various element contributions, the eutectic temperature, for

that alloy, as purchased, is 575.3°C. By looking at the

cooling curves, the primary phase growth temperature is almost

indistinguishable from the eutectic point, while the

undercooling is about 0.8°C. Micrographs show the sharp flake

structure of the silicon phase. By introducing a small amount

of magnesium (0.13%) the eutectic temperature is shifted to

573.8°C with 1.2°C undercooling. The associated microstructure

(Figure 3.12) now exhibits a slightly lamellar structure. A

further increas:e of the magnesium content resul ts in a

lowering of the eutectic temperature and a noticeable increase

in the undercooling. The microstructure, however, remains un-

changed, in that, the silicon has maintained its acicular

shape in sorne rE\gions while becoming lamellar in others.

When about 0.02% strontium was added to the alloy, the

eutectic temperature decreased to 571.5°C while the undercool-

ing remained constant. Thus, strontium additions lower the

eutectic temperature by about 3.S0C. The effect of the modifi-

cation on the microstructure is seen to be a complete change

of silicon phase morphology from acicular to very fine fibrous

63

(see Figure 3.13). The presence of an AlsFeSi phase and an

Al1s(MnFe) )Si2 phase now is more pronounced. Increasing the

magnesium content from 0.16% to 1.4% lowered the eutectic

growth temperature considerably. As observed in the synthetic

alloys, magnesium deteriorates the quality of the strontium

modification.

'\

{ .

64

'rable 3.7 The effect of magnesium addition on 413.2 commercial alloy

Mg Sr TNl TN2 .:1TN TEl TE2 .1TE (%wt) (%Wt) ( OC) (OC) ( OC) ( OC) ( OC) (OC)

0.02 568.6 575.0 6.4 574.5 575.3 0.8

0.13 570.2 573.7 4.9 569.2 573.8 1.2

0.42 569.7 573.0 3.3 569.0 570.4 1.4

0.56 569.4 574.6 5.2 566.5 568.5 2.5

1.25 569.0 574.2 5.2 561.8 564.5 2.7

Table 3.8 The effect of magnesium in the presence of strontium on 413.2 commercial alloy

Mg Sr TN1 TN2 ~TN TEl TE2 ~TE

(%wt) (%wt) ( OC) (OC) ( OC) ( OC) ( OC) ( OC)

0.02 0.024 570.3 574.2 3.9 572.3 571. 5 0.8 0.16 0.020 567.8 573 5.2 571.2 570.7 0.5

0.40 0.025 570.5 575.3 4.8 571 568.2 2.8

0.67 0.023 568.5 573.4 4.9 568.2 565.3 2.9

1.40 0.024 568.5 574 5.5 563.6 561.0 2.7

Note Cooling rate observed in the experiment using ceramic mold is 2°C/sec. TN1, TN2 and 4TN are the minimum, growth and undercooling temperature respectively for primary aluminum. TEl, TE2, and 4TE are minimum, growth, and undercooling temperature respectively, for eutectic phase.

' .. -

a) 0.02% Mg

c) 0.42% Mg

Figure 3.12 The evolution of eutectic microstructure as a function of rnagnesium content in an unmodified 413.2 commercial alloy.

- ,. r

a) 0.02% Mg

c) 0.42% Mg

Figure 3.12 The evolution of eutectic microstructure as a function of magnesium content in an unmodified 413.2 commercial alloy.

-------Il ')

d) 0.56% Mg

a) 0.024% Sr + 0.02% Mg

c) 0.025% Sr + 0.40% Mg

Figure 3.13 The evolution of eutectic microstructure as a function of magnesium content in strontium modified 413.2 commercial Alloy.

Il •

La .. ,\100 l'm ... -. - :,hJ-..I.',Y;, ''i

b) o. 02% Sr + O. 16 % Mg

ct) 0.023% Sr + 0.67% Mg

, j

3.4 - Magnesium and strontium Effect on th. Microstructure of 356.2 and 357.2 Commercial Alloy.

To confirm the observed effect of magnesium on well

modified commercial alloy. One specimen of each of the 356.2

and 357.2 alloys was modified with 0.02% strontium. As

expected, the alloy with the higher magnesium content, Le. the

357.2 alloy (0.5%-0.6\ Mg), is less modified than the 356.2

alloy (0.3-0.4% Mg) (see Figure 3.14).

a) 356.2 Alloy

Figure 3.14 The eutectic microstructure of a modified 356.2 and 357.2 alloy.

67

b) 357.2 Alloy

68

3.5 - Ettect ot Maqnesium on the Euteetie Temperature.

The depression of the eutectic temperature as a functian

of the magnesium concentration, as observed in the cooling

curves, is presented in Figure 3.15. Experimentally, it was

found that a correlation exists between the magnesium content

and the eutectic temperature oi the Al-7%Si, Al-11%Si, Al-

13%si synthetic alloys and the 413.2 commercial alloy. These

correlations are listed in Table 3.9 as linear equations.

These relationships, particularly that for the 413.2

Table 3.9 Linear relationships between the eutectic temperature and the magnesium levels.

Alloy Sr Modifi- Linear Equation Remarks *)

cation wt% Mg

AI-7%Si No TE=579. 0-17.6 (%Mg) 0.00-1. 01 Yes TE=574. 9-14.6 (%Mg) 0.00-1.21

AI-11%Si No TE=576. 4-9.5 (%Mg) 0.00-1. 06 Yes TE=573. 2-8.5 (%Mg) 0.00-1. 20

Al-13%Si No TE=576. 0-6.9 (%Mg) 0.00-0.98 Yes TE=573. 8-6.5 (%Mg) 0.00-1. 40

413.2 No TE=574. 6-8.7 (%Mg) 0.00-1. 25 Yes T(!=571. 4-7.8 (%Mg) 0.00-1. 40

*) Indlcate the magneslum leveis to where the Ilnear equations are valid.

commercial ailoy, may be use fuI ta estimate the magnesium

content by lnonitoring the eutectic temperature. As a compari-

son, a sirnilar linear relationship between the eutectic

temperature and the rnagnesiurn content has aiso been reported

for A356. 0 alloy70 :

-~--~

•

l

580 i

- t-U

575 l:- _ Wlthout Strontium 0 -QI

~ '-.. - Wlth Strontium

:1 ; !S70 .. QI - -CI. e ~ !S6!S u ;: u 560 QI '.

15 W

555 0 0.2 0.4 0.6 0.8 1 12 1 4

Magnesium content (%Wt)

a) Al-7%Si synthetic Alloy

- 578 (J 0 - 576

Q) =:; Wlhout Strontium ~ 574 :::s Wlth strontium ... ca 572 ~ Q) ~ 570 E

568 Q) 1- -' (,) 566 .-... (,) 564 t.'. Q) ... :::s 562 W 0 0.2 0.4 0.6 0.8 1 1.2 1.4

Magnesium content (%Wt)

c) Al-ll%Si synthetic Alloy

Figure 3.15 The depression of the eutectic temperature as a function of magnesium content for unmodified and strontium modified Al-Si alloys.

~

'1 "'"

•

580

-0 = Without Strontium 0 -CIl ~ ... - With Strontium ~ .. ca 570 ... 8. E CIl 565 t-u Ü

560 CIl ~ W

555 0 0.2 0.4 0.6 0.8 1 1.2 1.4

Magnesium content (%Wt)

a) Al-7%si Synthetic Ailoy

- 578 0 0 576 -! o Wlhout Strontium

574 ~ :::J ' . Wlth strontium ..

ca 572 .. CD a. 570 '. E ,:,

CD 568 "

1- 6 Co) 566 .-.. u 564 CD .. :::J 562 W 0 0.2 0.4 0.6 0.8 1 1.2 1.4

Magnesium content (%Wt)

c) Al-l1%Si Synthetic Alloy

Figure 3.15 The depression of the eutectic temperature as a function of magnesium content for unmodified and strontium modified Al-si alloys .

ô o -

578 r------------------, 576

574.

572 ................ .

~ ....

570 . la.

568 • Without Strontium • With Strontium

566

• 564 ~-----L------______ ~ ______ L_ ____ ~

o 0.2 0.4 0.6 0.8 1

Magnesium content (%Wt)

b) Al-13%Si Synthetic Alloy

578 r-----------------------------------~

576

574

572 570

568 566

564

562

5600 0.2 0.4 0.6

• Without Strontium

À With Strontium

0.8 1 1.2 1.4

Magnesium Content (%Wt)

d) 413.2 commercial Alloy

1.6

69

70

Tu = 582°C - 13.3 (%wt Mg)

This equation is only valid for the unmodified alloy with

magnesium concentrations in the range of 0.25% to 0.45% .

The introduction of about 0.02-0.03 wt% strontium to the

alloys markedly depresses the eutectic temp~rature. This was

not observed when magnesium alone was added at an equivalent

concentration, thus suqqestinq that maqnesium does Dot play

the important role in the eutectic silicon qrowth mechanism,

that strontium does. The decr(:·:se in the eutectic temperature

at high magnesium levels (>0.2%) is no doubt due to a simple

phase diagram effect. It is convenient, therefore, to further

discuss the Al-Si-Mg ternary phase diagram.

A two dimensional visualization of a part of the complex

Al-Si-Mg ternary system can be seen in Figure 3.16. The liqui-

dus surface is projected on to the base of a rectangular Al-

Si-Mg system. This rectangle is actually transformed geome

trically from triangular representation of the Al-Si-Mg

system, so that now the X-axis represents the silicon content,

the Y-axis represents Magnesium content and the origin (0.0)

is aluminum. The contours indicate the lines of equal liquidus

temperature. The eutectic temperature of Al-Si-Mg alloys is

represented by a line formed by the intersection of the

liquidus surface (line AB). This eutectic valley goes down ta

the ternary eutectic point as the Magnesium level increases.

The ternary eutectic temperature (555°C) is the lowest

position in the system where four phases (Liq-Al-M92Si-Si)

?

...

71

coexist in equilibr ium. The other informative feature on

Figure 3.16 is the existence of a quasi-binary line (sometimes

called pseudo-binary line). This line expresses the

correspondence between an intermediate phase (MgzSi) and

aluminum. The formation of Al-M92Si is always prior to the

binary Al-Si eutectic transformation.

The eutectic temperature of an Al-si alloy for a given f.lg

concentration can roughly be estimated by following the

12 "- '80 ... ... &70

..... "0 --~,

.... ..... '~o .... ....

, ... --- 6·0 --- llClUIdus 1-' .... .... , , tolo ::t:, ,

..... ---U .... , --- 6'0 .... W .... " -- tolO

f" '- -3 .... ,

.... ---1

, , 600 .- - __ .... \

~ . ' ~, \ -------------, ~ .... ; '. v;

S .... .... ". \ \ \ ~ ~~ \ \ \

Z , , ~. ,

\ \

0 , , .§ 'A\.UWIHIU ... \ 4( , 0; ,

\ \ , " ~ "

, ~ 0

.... \ Ir 0 ....

~ .... \ .

~ .... \ ~ 2 \ \ ..... 0

, , -'b 0 ,

Il • , -0 , ,

0

0 7 3 • S , 1 , ~ 10 14 IS

SlliCON-WEICHT. %

Figure 3.16 Part of a ternary Al-Si-Mg phase diagram. The eutectic valley of Al-si binary alloy goes down as magnesium concentration increased (71).

i .

72

solidification of the alloy itself. For instance, in an Al-

7%Si-O%Mg alloy, aluminum freezes first as a primary phase,

and then the melt freezes by following the X-axis until point

A. This is the eutectic temperature for the binary alloy

(about 577.5°C). Now, ~onsider an alloy having a composition

at point D. This alloy (Al-7%Si-O.6%Mg) will solidify along

the curve DE perpendicular to the contours and will move down

to the eutectic valley at point E. This is the eutectic

ternperature for that particular alloy composition where Al

residual and Si-rich solid solution forro as in a binary

eutectic system. Finally, the alloy will reach point B where

the remaining liquid solidifies as a ternary eutectic. Similar

reasoninq can also be applied to the Al-9%Si-o. 6%Mg alloy

(point C) where eutectic solidification eventually occurs at

point F. Thus we can estimate the eutectic temperatures for

aIl alloys used in the experiments. These estimates are

displayedt in Table 3.10 and are compared with the eutectic

temperatures observed from the cooling curves.

The experimentally determined values in table 3.10 do

differ sornewhat from those predicted from the phase diagram.

The differences are, however, small (less than <1%) and are

no doubt due to inacurracies in the phase diagram itself which

was generated on the basis of higher magnesium levels than

those used here. The eutectic temperatures measured in this

work are probably the best availaible for the Al-Si-Mg system

at low magnesium levels.

Table 3.10 Comparison of the eutectic temperature obtained from experiments and those from phase diagram.

wt% si wt% Mg TEl Tm % Dev

7.15 0.0 578.5 577.5 0.17 7.05 0.06 578.0 576.0 0.34 6.85 0.36 573.3 572.5 0.14 6.45 0.44 571.4 571.6 0.03 6.42 1.01 560.9 562.0 0.19

11.12 0.0 576.3 577.5 0.20 11.00 0.08 575.4 576.2 0.13 10.85 0.34 573.4 574.5 0.19 10.75 0.62 571.1 574.0 0.50 10.69 1.06 566.6 572.2 0.98

13.15 0.0 575.5 577.5 0.34 12.96 0.09 575.2 574.5 0.12 13.00 0.55 573.5 573.3 0.03 12.84 0.81 571. 6 572.8 0.20 12.76 0.98 567.6 572.4 0.84

Note: TEl and TE2 are the eutectic temperature determined from experiment and estimated from phase diagram respecti vely. % Dev = [Absolute (Tu-Tm) ] /T11 X100%.

73

3.6 - Image Analysis Resulta.

preliminary Note.

74

Image analysis was employed on the AI-7%Si and the

Al-13%Si synthetic alloys. The technique is sensitive enough

to quantitatively characterize the eutectic microstructure.

The analysis of both alloys is thought to be useful for

describing the magnesium effect on the eutectic

microstructure. The video images from secondary electron

emission (figure 3.17) were digitized qiving binary images

wh~ch could be further analyzed (fiyUre 3.18). As mentioned

previously, (chapter II), three parameters were used to

quantify the microstructure. The first one, the perimeter, was

used to determine the size of the eutectic particles. The

last two parameters, the shape factor and the aspect ratio,

were used to evaluate the particle shapes. Thus, if a particle

is spherical, the value of both the shape factor and the

aspect ratio is equal to 1. Conversely, these values increase

as the particles becomes more plate-like. Hence, as aspect

ratio and shape factor diminish, the degree of circularity is

improved.

The number of particles analyzed was about 200 per

measurement per sample. Each sample was measured three times

using a different object of interest. The total number of

particles measured was about 500-600 per sample. The results

of the image analysis are presented in the graphs shown in

75

a) Unmodified AI-7%Si Alloy

{ 1.,. •. .. ~ ~. . .. r-- ~ .. , . • /~",' f.- -. /-. ~- '/t,'· '1·· , ., ' " ~ ~. -.1 -.. .' •• ~ - '. ' --. 1.-- .. . ~, •. ." .... ,.~, • y ., ....... \ •.• 1:. ~ • • _.,' • '-........ '-, ,_' ~. _, .,. _/' _,. __ • _ 4 ~... .'_ ., l'." .. - ./1 . ~ .. - ' • ,.. ...... , J" ...... - Ile • "'. .. .• '-.'- ~.. ~.- .. " ..... ,- ." , .... - . .....,~ -. . - -. ,,- .-... .. --, . . - . .... .., .. . '- . - ... --- \. ~',. .:~..-, "'-~~/I .. ~ ~,.. ,,~~ "/' .1,-: .--, "'J." " • ~ -.1 · w~ • •. ,1'" ~ ,. 't;, ~. ... ... ~'

~~~~I" ~ tJ~HJ 30KU ~ -~ ... -~~. 10~m W03 f. b) Modified AI-7%Si Alloy

Figure 3.17 Secondary electron images.

a) Unmodified Al-7%Si Alloy

b) Modified Al-7%Si Alloy

Figure 3.18 Binary images which are created from video/secondary electron images.

77

Figures 3.19 to 3.24. Both the AI-7%Si and the AI-13%Si alloy

clearly show the alteration of silicon eutectic particles, (as

characterized by their shape and dimension), dS a function of

either magnesiuiTI or strontium content. On each graph, two

different behav~ours are evident in the variation of the

microstructural pal"'ameters wi th and wi thout strontium.

Without strontium, the mean value of the perimeter, the shape

factor and the aspect ratio of the eutectic phase, as weIl as

their standard deviation, fluctua te very w~ùely. Thus the

silicon phase is very non-uniform in the absence of strontium.

In the presence of 0.02% strontium, as magnesium increased,

the mean values increased slowly and their deviation steadily

broadened. In general, the presence of magnesium tends to

produce a non--uniform eutectic structure while strontium

homogenizes the microstructure.

Hypoeutectic Synthetic Alloy

ln the pure AI-7%Si alloy (see Figures 3.19 to 3.21 ), the

average perimeter of the silicon phase is about 48}.J.m. This

value fluctuates with increasing magnesium content (0.36-0.44

wt%) as does the deviation of the mean value of the perimeter.

This implies that a fine silicon phase is present in the

structure as indicated earlier in the micrographs. The shape

factor, regardless of the magnesium content, remains

essentially constant. The aspect ratio, however, steadily

decreases from 3.6 to 2.7 as the m""\gnesium content increases.

78

In the absence of strontium, low concentrations of magnesium

change the eutectic shape leading to a very non-uniform

microstructure (reflected by the broadening of deviations) .

In the presence of strontium, the degree of modif ication

is appreciable. Tne eutectic silicon phase i3 fibrous with ù

high degree of circularity. without magnesium, the average

value of the perimeter is 7 J.Lm with no significant deviation.

The shape factor is about 1.7 and the aspect ratio is about 2.

Introducing small amounts of magnesium to the alloy causes

both the size and the shape of the eutectic phase to increase.

This is reflected, not only by the mean values but aiso by

their deviation. Thus the microstructure becomes more hetero-

geneous as magnesium increases. Finally, with about 1.2%

magnesium content, the average of the eutectic perimeter is

about 30 J.Lm; the shape factor is about 4, and the aspect ratio

is 2.4 approaching that of the unmodified alloy. This alloy

(as shown in the previous micrographs) is characterized by a

mixture of fibrous, lamellar and acicular silicon. It can be

concluded that magnesium does

modification.

indeed impair strontium

,'1

'\

{ ..

19

1 2 0 r 1 1O • Wlthout Strontium 100 Â Wlth Strontium , - 90

1 E 80 , .:, ~ 70 1 T CI> J , ,

Q) 60 1

E 50 • 't: 1 CI> 40 l r Q, 1

30 • 4- + ,

20 10 l

0 1 1 1 1 ! 0,0 0,2 04 06 0.8 1 0 1 2 ~ 4-

Magnesium Content (%Mg)

Fig 3.19 The perimeter of silicon particles as a function of magnesiurn concentration (Al-7%Si alloys).

10~----------------------------------

• Without Strontium 8 • Wlth Strontium

t f 2

O+---~~---+----~----~-----rl----~I----~ 0.0 0.2 0.4 0.6 0.8 1 0 1 2 1.4

Magnesium Content (%Wt)

Fig 3.20 The shape factor of silicon particles as a function of magnesium concentration (Al-7%Si alloys).

T

'. 1-.

... 1 f

t ï (

l

51 1

4

0 :;: ca a: 3 T - • (,,) CI)

~

C-I/)

ct

: j. 00

1

02

..

• Without Strontium .. Wlth Strontium

•

. ,

Magnesium Content (%Wt)

Fig 3.21 The aspect ratio of silicon particles as a function of magnesium concentration (Al-7%Si alloys).

Eutectic synthetic Alloy

80

As shown in the micrographs in Figure 3.10 the pure Al-

13%5i alloy is characterized by sharp-flakes of the eutectic

silicon particles distributed throughout the aluminum matrix.

Image quantification revealed that the perimeter of the

silicon phase is quite variable (ranging from 20 to 80 ~m),

with inconsistencies in the standard deviation (see Figures

3.22 to 3.24). In fact, this reflec:ts the nature of the

eutectic silicon when magnesium is added. The same tendency

was also observed for the shape factor, espe~ially its

standard deviation, suggesting that the silicon phase becomes

non-homogeneous. The aspect ratio, however, decreased fram

4.5, in the absence of magnesium, to about 3.0 for a 0.98%

magnesium content. 'ro summarize, by increasing magnesium

concentration, the silicon shape is altered from sharp

".1 !,.

81

acicular flakes to rough-faceted larnellae, and the size

increases s l ightly in sorne parts and coarsens in others,

especially with about 0.98 wt% Mg.

As described previously, in the presence of strontiu~,

the eutectic silicon phase is fibrous and very finely divided.

This phase has a mean value of the perimeter of 10 ~m, shape

factor of 1.8, and aspect ratio of 1.9. All parameters exhibit

very small deviations indicating a homogeneous microstructure.

The addition of rnagnesiurn to the alloy steadily increases not

only the rnean value of these pararneters but their deviatjon as

well. Particularly in the case of the 1.4% Mg addition, the

silicon size is almost doubled from 10 to 19 ~rn.

100

80 1 - f E

1 l a, 60 '- T CD

1 1$ i E 40 i 1: T CD 4 Q.

+ 20 1 • 4

0 1 1 1 1

00 02 04- 06 08 1 0 1 2 1 4- 1 6

Magnesium Content (%Wt)

• Wlthout Strontium ... With Strontium

Fig 3.22 The perirneter of silicon particles as a function of magnesium concentration (AI-IJ%Si alloys).

• • 'c f 1

~ ~ (

\'

~ " ~ \ • 1 l -~ ~ • i .'

1

82

20 i

1 • Wlthout Strontium

~ 15

1 0 .. u ra

10 f IL Q) Q.

l 1 ra J: fi)

5I A

• Wlth Strontium

• • T .. • , •

01 ~- -+-. - -

00 o 2 .] -1- o t3 J S i J ~

~

Magnesium Content (%Wt)

Fig 3.23 The shape factor of silicon particles as a function of magnesium concentration (Al-13%Si

.2 .. ra a: .. u Q) Q. II)

<r:

alloys) .

60

5.0

4.0

T 30 , 2.0 • 1 a

• ..

• Wlthout Strontium

À Wlth Strontium

i

.. . 00 +---+---...-----+---~--r-------~- ---.

0.0 02 :J 4 J6 08 1 J 2

Magnesium Content (%Wt)

Fig 3.24 The aspect ratio of silicon as a function of magnesium concentration (Al-13%Si alloys).

83

3.7 - Aspects To Es Considered

Two questions arise from the experimental results

reviewed in the preceding section. First of all, in the

absence of strontium, why and how does maqnesium change the

silicon phase, and secondly, why does magnesium mask the

strontium modification ?

1. Magnesium alteration of the eutectic silicon phase.

Magnesium, up to now has been reported to have various

effects in these alloys. Sorne researchers have reported that,

magnesium refines the unmodified alloy 3.69 while others have

observed that magnesium coarsens the eutectic micro­

structura24• The role of ma~'nesium on the nucleation and growth

mechanism of eutectic silicon is far from clear. This

confusion as to the effect on the eutectic structure is

perhaps understandable since magnesium is added to foundry

alloy for purposes of heat treatment, and so little attention

has been paid to its effect on the eutectic microstructure.

From our experiments, it is hypothesized that magnesium

changes the silicon phase by two possible mechanisms.

The first is by influencing the silicon growth mechanism, (in

a different way than strontium does), so that the silicon

crystals grow in many directions and not just in the favoured

crystallographic direction (unlike strontium induced

twinning). The rounded appearance of the silicon particles is

due to the higher frequency of high angle branching during

r

.~

----------------------------~~, .......... , ...... .

84

crystal growth. The reason why magnesl.urn is capable of

affecting the silicon growth mechanisrn is due to the fact that

the ratio of its atornic radius to that of silicon approaches

the w\lue of 1.65, a value where modif ication is thought ta be

rnost effecti ve55• The second hypothesis is based on the

accumulation of magnesium atorns ahead of the eutectic silicon

phases. This causes a shift of their melting points down

extensions of the liquidus surface resulting in a decrease in

freezing rate leading to coarsening of th8 structure.

2. Effe~t of maqnesium on strontium modification.

The most reasonable approach to answer the question of

why magnesium masks the modification effect is that rnagnesiurn

reacts with strontium to form a complex intermetallic

compound. Such compounds may form in the liquid before the

eutectic reaction. Investigation by electron rnicrascopy uSlng

EDX enabled the identif ication of compounds that could be

responsible for modification degradation (see Figure 3.25).

The chemical stoichiometry of the compounds was finally

determined using WDX from an electron microprabe by comparing

the X-ray intensity emitted by the elements of interest. It

was found that the following complex compound is formed:

Mg2Sr (Si3Al4 ) •

These compounds appear as very bright spots on the back -

85

scattered electron ~rnage on the CTR (see Figure 3.26). Their

formation is very cornplex in the Al-Si-Mg-Sr system and 50 far

no information on thern appears in the literature.

Figure 3.25 Intermetallic compound found in the Al-Si-Mg-Sr system (Mag 5000 X)

86

Figure 3.26 The complex Mg2Sr(Si1Al~) intermetallic compound (bright spot). By increasing the image contra st (the lower micrograph), the compound is more pronounced and can be clearly distinguished from the eutectic silicon phase (Mag 4500X).

87

3. Some other important factors to be discussed

In real foundry practice, magnesium, in the range of 0.3

to 0.6 wt%, is commonly added to Al-Si alloys in order ta make

heat-treatable alloys. Beyond 0.7 wt%, as reported by several

researchers mentioneù in Chapter l, there is no further

improvement in the mechanical properties. This perhaps is due

to the alteration of the microstructure, so that the strength-

ening effect of Mg2Si particles may be compensated by the

coarsening effect of the silicon phase. In the case of

strontium modified alloys, there is no information in the

literature regarding the influence of magnesium on the

loicrostructure. Based on the current exper iments, however,

(see Figures 3.4 through 1.13) alloys with magnesium levels

between l.3 0.6 wt% exhibit microstructures which are

typically dominated by the fibrous silicon phase. Hence, there

may be no significant effect on the mechanical properties. On

the other hand, increasing the magnesium content in excess of

0.6 wt% affects the modification markedly, so that the

microstructure is no longer homogeneous but consists of a

mixture of the fibrous, lamellar, or even acicular structures.

The higher the magnesium concentration, the more pronounced

the damage to the modification. As a result, beyond the

optimum magnesium content, one may expect a deterioration of

the mechanical properties. Therefore, it is not recommended to

add magnesium in excess of the 0.6 wt% level to strontium

modified Al-Si alloys.

88

Recalling the silicon phase characterization (see Figures

3.19 through 3.24), unmodified alloys exhibit fluctuations of

the silicon size and shape parameters, as weIl as fluctuations

in their standard deviation. This, perhaps, is not a 'true'

standard deviation, which according to statistical principles

is defined as a tool to assess the accuracy of the rneasure­

ments or to assess the significance of the error introduced

dur inq repeated measurements under the same exper imenta 1

conditions. Conversely, in the case of the present micro­

structural characterization, the standard deviation should be

considered solely as a tool to assess the homogeneity of the

silicon phase. As observed previously, particularly in the

strontium modified alloys, the higher the rnagnesium level the

wider the standard deviation. The standard deviation, in this

particular case, reflects the nature of the microstructure.

,

( ,

89

Chapter IV

CON C LUS ION S

1. In Al-7%Si, Al-11%Si, and Al-13%Si synthetic alloys,

a low concentration of magnesium changes the silicon phase.

In general, i: was found that Magnesium neither refines

nor coarsens the eutectic phase, but definitely reduces

the degree of homogeneity of the microstructure. At high

levels (eg.1 wt%), however, the magnesium slightly refines

the silicon phase.

2. The effect of magnesium in Al-13%si eutectic alloys is

most distinct. It changes the eutectic silicon phase

from sharp-acicular flakes to rough faceted lamellae.

3. Magnesium has a negative effect on strontium modification,

such that its presence changes the microstructure from

a fully modified to a partially modified one.

4. This effect is most likely due to the formation of a

complex Mg2srAl~Si3 intermetallic compound which probably

is formed prior to the eutectic transformation.

5. Magnesium, unlike strontium, does net play an important

role in the eutectic silicon solidification.

6. The depression of the eutectic temperature in Al-Si alloys

as magnesium increased is due to phase diagram effect.

7. For feundry applications, it is suggested that magnesium

not be added in quantities higher than O.6%Wt in strontium

modified Al-si casting alloys.

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90

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r i ,

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~-

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f

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93

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97

Appendix

Chemical composition of the Alloys

Table 1. commercial Alloys as Purchased

Alloy si Fe Cu Mn Mg Zn wt% wt% Wt% wt% wt% wt%

356.2 6.03 0.430 0.034 0.010 0.290 0.020

357.2 6.52 0.224 0.038 0.120 0.586 0.022

413.2 10.55 1.280 0.010 O. 020 0.020 0.010

Table 2. Al-7%Si Synthetic Alloys

Si Fe Cu Mn Mg Zn Sr wt% wt% wt% wt% wt% wt% wt%

7.15 0.004 0.003 0.001 0.00 0.001 -7.05 0.003 0.004 0.003 0.06 0.002 -6.85 0.004 0.003 0.004 0.36 0.003 -6.45 0.002 0.002 0.003 0.44 0.002 -6.42 0.003 0.003 0.002 1.01 0.003 -7.10 0.004 0.002 0.002 0.00 0.001 0.030 7.00 0.003 0.002 0.001 0.08 0.002 0.030 6.78 0.002 0.003 0.002 0.55 0.002 0.028 6.67 0.003 0.001 0.003 0.81 0.001 0.029 6.56 0.001 0.004 0.004 0.98 0.001 0.025

{

98

Table 3. Al-l1%Si Synthetic hlloys

Si Fe Cu Mn Mg Zn Sr wt% wt% wt% wt% wt% wt% \-Jt%

11.12 0.004 0.002 0.001 0.00 0.001 -11.00 0.002 0.001 0.003 0.08 0.002 -10.85 0.001 0.003 0.001 0.34 0.001 -10.75 1 0.002 0.002 0.003 0.62 0.002 -10.69 0.002 0.003 0.002 1.06 0.003 -

11.10 0.002 0.002 0.002 0.00 0.001 0.021 11.05 0.003 0.004 0.001 0.08 0.002 0.018 10.93 0.002 0.003 0.002 0.33 0.002 0.020

10.67 0.004 0.001 0.001 0.60 0.001 0.025

10.46 0.001 0.004 0.003 1.20 0.001 0.021

Table 4. Al-13%Si Synthetic Alloys

si Fe Cu Mn Mg Zn Sr

wt% Wt% wt% wt% wt% wt% Wt% -

13.15 0.003 0.002 0.001 0.00 0.001 -12.96 0.002 0.001 0.002 0.09 0.002 -13.00 0.001 0.003 0.001 0.55 0.001 -12.84 0.003 0.002 0.003 0.81 0.001 -

~

}4. 12.76 0.002 0.002 0.001 0.98 0.002 -, ;:

l ~ ~

13.10 0.001 0.002 0.003 O.~Q 0.001 0.024

13.05 0.003 0.001 0.002 0.08 0.001 0.030

t, t

t 12.93 0.002 0.003 0.001 0.57 0.002 0.029

12.98 0.003 0.001 0.001 0.70 0.001 0.030

12.57 0.001 0.003 0.003 1.40 0.003 0.034 , t "

t f -"lI-

~ <

~:

99

\

Table 5. 413.2 Commercial Alloys

Si Fe Cu Mn Mg Zn Sr Wt% wt% wt% wt% Wt% wt% wt%

10.55 1. 28 0.01 0.02 0.02 0.01 -10.96 1. 30 0.02 0.02 0.13 0.02 -10.67 1.17 0.01 0.01 0 .. 42 0.01 -10.35 1. 25 0.02 0.03 0.56 0.01 -9.87 1.19 0.01 0.01 1.25 0.02 -

10.55 1. 38 0.01 0.02 0.02 0.1"11 0.024 10.75 1. 25 0.01 0.02 0.16 0.01 0.020 10.45 1. 35 0.02 0.01 0.40 0.02 0.025 10.10 1. 23 0.01 0.01 0.67 0.01 0.023 9.76 1. 34 0.01 0.03 1.40 0.01 0.024