effects of magnesium, silicon, and …...effects of magnesium, silicon, and strontium on the...
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EFFECTS OF MAGNESIUM, SILICON, AND STRONTIUM ON THE OXIDATIQN OF MOLTEN
ALUMINUM
BY
KEITH DENNIS
Department of Mining and Metallzirgy
McGill University, Montreal
June 1999
A Thesis submitted to the Faculty of Graduate Studies and Research
in partial fulfillment of the requirements of the degree of
Master of Engineering
O Keith Demis, 1999
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Oxidation experiments were carried out on six different alurninum alloy
compositions in order to determine the effects that silicon, strontium, and magnesium
have on the oxidation behavior at temperatures between 700 and 800°C. This was
determined through weight gain studies using a thermogravirnetric balance, microscopie
examination of sample surfaces and cross-sections, and chemical analysis of the phases
formed during oxidation.
The addition of 12% silicon to aiuminum prornoted the formation of a-alumina
nodules on the sample surface at 800°C by reducing the viscosity of the molten metal to
allow for better feeding of fiesh metal through cracks in the top oxide. Strontium was
found to prevent the formation of oxide nodules in an Ai-12%Si ailoy through the
formation of a strontium containing oxide. Similarly, in a commercial 356 alloy,
strontium greatly reduced the extent of oxidation by preventirig the preferential oxidation
of magnesium. It is believed that the mechanism responsible for this is also through the
formation of a strontium containing oxide.
Oxidation experiments conducted on an Al-3%Mg alloy revealed that magnesium
preferentially oxidizes to form two oxides, firstly MgO, and secondly MgA1204, which
forms only from the reaction of Mg0 and aluminurn. Weight gains for this alloy were
greatly affected by the oxidizing temperature. Higher temperatures shortened the
incubation time to breakaway oxidation and decreased the tirne required to oxidize al1 the
magnesium fiom the alloy.
A h de déterminer les effets qu'ont le silicium, le strontium et le magnésium sur
le comportement en oxydation à des températures comprises entre 700 et 800°C, des
s X ~ & & m f i ~ pb,.,.,b,s ont ité f i tes sur six diffé:tr,!es compositions d'alliages d 'a lu i i i iun
L'oxydation a été évaluée par des études de gain en poids B l'aide d'une balance
thennogravimétrique, par des examens microscopiques des surfaces et des tranches des
Cchantillcns, et finalement, par des analyses chimiques des phases formées pendant
l'oxydation.
L'addition de 12% de silicium à L'aluminium favorise la formation de nodules
d'alumine alpha sur la surface de l'échantillon à 800°C en réduisant la viscosité du métal
liquide permettant ainsi une meilleure alimentation en métal fiais au travers des fissures
de la couche d'oxyde. Il a été trouvé que le strontium empêchait la formation de ces
nodules d'oxyde dans l'alliage AI-IZ%Si en formant un oxyde contenant du strontium.
D'une manière similaire, dans un alliage commercial 356, le strontium réduit
considérablement l'étendue de l'oxydation en empêchant l'oxydation préférentielle du
magnésium. On pense que le mécanisme responsable de ceci est aussi relié à la formation
d'un oxyde contenant du strontium.
Les expériences d'oxydation gour un alliage AI-3%Mg ont démontré que le
magnésium s'oxyde préférentiellement pour former deux oxydes, d'abord MgO, puis
MgAl2O4 qui se forme seulement par la réaction de Mg0 avec l'aluminium. Les gains en
poids pour cet alliage ont été considérablement affectés par la température oxydante. Des
températures plus élevées ont raccourci le temps d'incubation avant qu'une oxydation
puisse être observée et ont diminué le temps nécessaire pour oxyder tout le magnésium
contenu dans l'alliage.
iii
ACKNOWLEDGEMENTS
1 wouid like to express my sincere gratitude to Prof. R.A.L. Drew and Prof. J.
Gruzieski for their supervision and guidance throughout the entire project. Their
continueri vitdity md hterest h the pr~ject w:e encouraghg exmples :O follow.
financial support was also greatly appreciated.
1 wouid like to th& d l the lab technicians who made it so much easier to
perforrn al1 the different tests included in this work. Speciai thanks go to Helen Campbell
for help with the SEM, Monique Riendeau for help with x-ray difiaction, Ray Langlois
and Edwin Femandez for their tirne patience and generosity and fmally Glenn Poirier for
assistance with the Electron Probe Microscope. 1 would also like to thank Florence Paray
for her wisdom and expertise in translation.
1 wish to thank ail my fellow graduate students who made my stay at McGill
filled with fun, excitement, and leaming. To Dave, Neil, Sandra, Stacey, and Tara, thank
you al1 so much for the amazing mernories and laughs in the last six years. Special thanks
go to the ceramic group who had to put up with me day in and day out in the laboratory.
You have al1 made my time of research here at McGill a mernorable one.
Finally, 1 wish to thank my parents for their love and support (financial and
otherwise). The rnajoriry of this work beiongs to them because, without their help, none
of this would have been possible. Additional thanks to Byron and Shelley for their
unfettered support.
TABLE OF CONTENTS
. . ........................................................................ ABSTRACT II . . RESUME ....................m..m..m...m.............m.............m~.m............ ..III
...................................................... ACKNOWLEDGEMENTS iv TABLE ........................................................ CONTENTS v
. . LIST OF FIGURES ...........................m..........................m........ VII
LIST OF TABLES . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . xi
NTRO
2.1 Oxidation .................................................................................. 3
2.2 Oxide Formation ........................................................................ 4 ................................................................ 2.2.1 The Pilling-Bedworth Ratio 4
2.3 Oxidation Kinetics ...................................................................... 6 ........................................................................ 2.3.1 Thin Films and Scaies 7
2.3.1.1 Thin Films ................................................................................. 7 2.3.1.2 Scales ....................................................................................... 9
............................................................... 2.3.2 Dependence on Temperature i O ................................................................................. 2.3.3 Alloy Oxidation i r
2.4 Oxidation of Pure Aiuminum ........................................................ 12 ................................................................................. 2.4.1 Initial Oxidation 13
...................................................................................... 2.4.2 Scale Growth 14
............................................ 2.5 Aluminum-Magnesium AUoy Oxidation 16 . . ........................................................................ 2.5.1 Magnesium Oxidation 17 ........................................................ 2.5.2 Alumhum-Magnesium Oxidation 18
2.5.2.1 Oxidation in the Solid State ......................................................... 19 ................................................................ 2.5.2.2 Liquid State Oxidation 22
2.6 Effect of Silicon on Oxidation ...........m............................................ .27
................................................ 2.7 Effect of Strontium on Oxidation m o m o a 2 8
2.8 Summary .................................................................................. 29
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3 . EXPERIMENTAL PROCEDURE 30 3.1 h w Materials ........................................................................... 30
3.2 Melting and Casting ..........................................m......................... 31
3.3 Sample Preparation ..................................................................... 32
3.4 Crucible Fabrication .................................................................... 32 ................................................................. 3.5 Oxida tion Esperiments 33
......................................................................... 3.6 Sample Analysis 36 . . ....................................................................... 3.6.1 Emission Specîroscopy 36 ............................................................................ 3.6.2 Opticai Microscopy 36
............................................................ 3.6.3 Scanning Electron Microscopy 37 ............................................................................... 3.6.4 X-Ray Diffraction 37
..................................................................... 4.1 Alloy Compositions 38
4.2 Pure Aluminum .......................................................................... 39
.................................................. 4.4 Synthetic 413 Alloy Containing Sr 43
5.1 Oxidation of Pure Aluminum ......................................................... 61 .............................................. ..... 5.2 Effects of Silicon ,, .................. .. 62
....................................... 5.3 Effects of Strontium on Synthetic 413 Alloy 63 ............................................................... 5.3.1 X-Ray Difiaction Analysis 65
.................................................................................... 5.3.2 Microanalysis 66
.......................................... 5.4 Oxidation of Synthetic 5000 Series Alloy 66 ................................................................................. 5.4.1 Initial Oxidation 67 ................................................................................. 5 .4.2 Rapid Oxidation 70
................................................. 5.5 Oxidation of Commercial 356 Alloy 74
................................... 5.6 Effects of Strontium on Commercial 356 Alloy 76
TABLE OF FIGURES Figure 2.1
Figure 2.2
Figure 2.3
Figure 2.4
Figure 2.5
Figure 2.6
Figure 2.7
Figure 2.8
Figure 2.9
Figure 2.10
Figure 3.1
Figure 3.2
Figure 3.3
Figure 3.4
Figure 4.1
Figure 4.2
Figure: 4 3
Figure 4.4
Oxide stresses according to Piliing-Bedworth ratios .................. 6
General classification of weight gain kinetics dunng oxidation ..... 7
Schematic of the diffusion occurring in the Cu-Cu20-O2 system .... 10
Mode1 for the solid-state breakaway oxidation ......................... 22
Stability diagram for MgO. spinel and A1203 ................ .. ..... .. 24
Selected data indicating the dependence of the oxidation rate on ............ temperature and magnesium content of a ACMg-Si alloy 25
...................... Effect of Po2 on the gnvoth rate of a AI-Mg alloy 26
Effect of silicon additions on the incubation period of a Al-Mg Alloy ............................................................................. 27
Solubility of oxygen in aluminurn with silicon additions .......... .... 28
............. SEM pictures of top oxide formed at 740°C for 24 hours 29
Graphite mold used for casting of ailoys ................................. 32
Schematic illustration of the drain-casting process, (a) fil1 mold with slip, (b) liquid extracted through mold leaving compact layer along walls, (c) excess slip drained, and (d) casting removed after partial drying ................................................ 34
......... Experimental set-up to measure the oxidation weight gains 35
Location of samples taken from each alloy for chemical analysis .. -36 Weight gain curves for pure aluminum at 700 and 800°C ............ 39
......... Surface oxide of pure aluminum sample oxidized at 750°C 40
X-ray diffraction pattern of a pure aluminum sample surface oxidized at 800°C .................................................... 40
Weight gain curves for synthetic 413 ailoy a t 700,750, 800°C ...... 41
vii
Figure 4.5 Oride surfaces of the synthetic 413 ailoy ................................ 42
Figure 4.6
Figure 4.7
Figure 4.8
Figure 4.9
Figure 4.10
Figure 4.1 1
Figure 4.12
Figure 4.13
Figure 4.14
Figure 4.15
Figure 4.16
Figure 4.17
Figure 4.18
Figure 4.19
Figure 4.20
X-ray diffraction pattern of 413 mock aUoy oxidized at 750°C for 45 houn. ..... .. . .. . . .. . . . . ..... .. .... . ..... .. . . .. ... .. .... .. . . .. ... ... ... . . .42
Weight gain curves for the synthetic 413 ailoy containing 350 ppm Sr a t 700,750,800°C ................................... A Micrographs showing the effect of Sr as a modifler in the 413 mock slloy .............. ................oo...~.......................... 44
Microscopy of the oxide surface of the synthetic 413 alloy containing 350 ppm Sr oxidized at 800°C. .. .. ... . . ... .. . .. .. .. .... .. .... 45
X-ray diffraction pattern of synthetic 413 alloy containing 350 ppm Sr oxidized at 750°C for 45 hours .............................. 45
Weight gain curves of the commercial 356 alloy at three different temperatures.. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 46
Low magnification of top oxide surfaces and of the commercial 356 alloy ... ..... . ..... . ... ...... .. . . . . . . ... . .. . . .. . .. . . . . . .. . . 4 7
High magnification micrograph of the surface oride on commercial 356 alloy samples.. . . . . . . . . . . . . . . . . . . . . . ... . . . . . . . . . . . . . . . . . . 47
X-ray diffraction pattern of the commercial 356 alloy oxidized a t 750°C.. .... . . . . . .. .. .. . ...... . . . .. .. ... . . . . .. . . . . . 4 8
Weight gain curves for the commercial 356 Alloy containing 250 pprn Sr at three different temperatures ....... .... ..... .. ...... 49 Micrographs showing the modification of the silicon phase in a commercial 356 alloy due to the presence of 250 ppm Sr.. . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . .. 49
Low magnification of commercial 356 alioy containing 250 ppm Sr oxidized at 750°C... .. ..... .... . .. .... .. . ..... . . .. ..... .. ....50
Coherent surface oxides formed on commercial 356 alloy containing 250 ppm Sr.. . .. . . . .. .. .... . . .. .. . . .. .. . . . .. ... ... ... ... . ..... . . ... 50
Localized oxide eruption (700°C) .. .... .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 51
X-ray diffmction pattern of commercial 356 alloy containing 250 ppm Sr ...................................... . . . . ........... . . .. 51
viii
eb Figure 4.21 Specific weight gains for the synthetic 5000 series ailoy at three dinerent temperatures. . . . . . . . . . . . . . . . . . . . . . . . . . a a . . . . . . a . . . . . . . ..53
Figure 4.22 Specific weight gain of the commercial 5000 series alloy at 700°C.. 53
Figure 4.23 Micrographs showing caulinower-üke morphology of surface oxide of the synthetic 5000 series alloy at '150°C... . . . . . . . . . . . . . . . . . . . . .55
Figure 4.25 Spikes penetrating the top oxide surface. (800°C, 3 hours). . . . . . . . . .. 56
Figure 4.26 Micrograph showing the spinel and al.uminum phases present in an oxide cluster oxidized for 45 hours.. . . . . . . . . . . . . a . . a . . . a . . . . . . . . .. . 57
Figure 4.27 Elemental maps showing the distribution of Al, Mg, and O throughout an oxide cluster a) secondary electron image of entire cluster, b-d) elemental maps of a portion of cluster (750°C, 22 hrs oxidation). ................................................... 58
Figure 4.28 X-ray diffraction patterns for the synthetic 5000 series alloy ........ 60
Figure 5.1 Cornparison of weight gains of pure aluminum and 413 alloy at 800°C.......... .... mma.aa...m.~..a...a...maaa..aaa.aa.aaa~aa.aCaaaa.a..a~~..aa.62
Figure 5.2 Cornparison of the initial oxidation behavior of pure aluminum, 413, and 413 containing strontium ......... .a.m.m.. . . . . . . . . . .amma.a.a. .~~.~ 64
Figure 5.3 Cornparison of x-ray diffraction peaks of 413 and 413 containing strontium alloys ............. aa....aa.ama.aaaaam..a.m..mm..aa...a....a.m.a.aaa.a65
Figure 5.4 Weight gain for the synthetic 5000 series alloy at 700°C showing three different oxidation kinetics.. .. .. .. . ..... . ... . .. . . . . . . . .. . . . .. . . .. .. ..68
Figure 5.5 An example of the initial oxidation period of a 750 and 800°C sample ............. a.m...a.a.oaa.aao.aa.8..aaa.a.a....oam..aa.aaa.aa~..ma.aaaaaaaa69
Figure 5.6 Elernental scans of a) Al and b) Mg of a selected part of a sample oxidized at 8D0°C ....+............................................... 70
Figure 5.7 The presence of small crystals (potentially MgO) on the sample surface near the point of breakaway initiation .........~.. .... 72
Figure 5.8 Schematic of the proposed mechanism responsible for the initiation of breakaway oxidation. a) Formation of magnesium vapour pockets below the top oxide layer, b) Cracking of the oxide layer along the sides of the sample due to spinel crystal formation. c) Flow of metal out of crack and release and immediate oxidation of mapesium vapour pockets.. ................. 73
Figure 5.9 Specific weight gains of the commercial 356 alioy ..................... -75
Figure 5.10 Cornparison of the initial oUdation peRods of the commercial ..................................................................... 356 alioys.. -77
Figure 5.11 X-ray diffraction pattern showing small peaks matching strontium-containing phases. .............................................. -78
LIST OF TABLES Table 2.1. P-B Ratios of some metais ........................................................ 5
Table 2.2: The effect of temperature on the oxidation behavior of various .................... ....................................................... metals ...- I l
Table 2.3. Various Gibb's free energies for the formation of spinel .................. 18
Table 3.1. Chernical analysis of the raw materials used ................................ -31
........ Table 4.1. Composition (in weight %) of the various aluminum alloys used 38
Table 5.1. Average induction periods ....................................................... 67
1 INTRODUCTION Throughout the processing of aluminum, there exist many stages where aiurninum
is exposed to air. For instance, during many melt transfer operations, the molten metal is
exposed for brief intervais. Additionally, many commercial melt holding furnaces are
open to the atmosphere. Due to alurninurn's high finity for oxygen, oxidation on a
fieshly exposed metai surface will occur rapidly. The concept of preventing aiurninum
oxidation throughout the industrial process is unredistic and therefore the concem of
controlling that oxidation is of utmost importance. With new treatments, ailoy
compositions, and other variables being introduced, the oxidation phenornenon has
become much more complicated.
Oxidation is dependent on the chemistry of the metal. With respect to aluminurn
alloys, common alloying elements such as magnesium and lithium have been found to
oxidize in an accelerated manner. This impact is felt in the overall process efficiency.
More and more aluminurn industries are becorning concemed with the efficiency of their
process and oxidation is one of the major problems. The majority of metal loss is due to
oxidation, whether through direct reaction with oxygen or through metal entrapment in
the oxide slag. This slag, or 'dross', contains a mixture of oxides, sdts, and retained
-
metal, which is typically treated for recovery. This process, however, can be lengthy,
expensive and at times, inefficient.
In order to prevent or decrease the amount of oxidation, it is first necessary to
understand the mechanisms and reactions involved. The objective of this work is to
examine the oxidation behavior of various synthetic and commercial aluminum alloys in
order to observe the eFfects of magnesium, silicon and strontium at cornon foundry
temperatures. This is carried out through the use of a thermogravirnetric balance and
rnicroscopic examinatioo of the oxidized samples.
2. LITERATURE REVIEW
2.1 OXIDATION Oxidation can be defined as the removal of one or more electrons fkom an ion,
atom, or molecule. In the case of aluminum, the electrons are lost by the aluminum atoms
and are taken up by the oxygen atoms to form their respective anions through the following
reactions:
2 Al -+ 2 ~ l ) ' + 6e- (2- 1)
312 O2 + 6e- + 302- (2.2)
The respective ions then combine to form the oxide, aiumina:
2 AI^+ + 302- -, &O3 e3) In order for oxidation to take place, the ions of both oxygen and duminum must
come into contact with each other. Initially, this condition is met since in most cases
a1uminu.m will come into direct contact with air. The continuation of üiis reaction however,
is limited by the formation of the initial oxide. If a complete oxide forms over the surface,
m e r oxidation can only occur if the ions can somehow diffuse through this preliminary
oxide layer. If this is the case, as the oxide grows thicker, it will take longer for the ions to
diffuse through, thereby slowing down the oxidation. Since oxidation is often diffusion
Chapier 2 Li'iera f ure Review 4
dependent, factors such as temperature, and protectiveness of the oxide become important
considerations for oxidation studies. These parameten will be dealt with later.
2.2 OXIDE FORMATION
Virtually no metal cemains unafTected by a gaseous environment. The problem
therefore reduces to a consideration of the extent to which coatings fomed by reaction
with the gas are protective. The protectiveness of an oxide is a term used to describe how
well the initial oxide layer that forms protects the metal substrate fiom further oxidation.
Essentially three variations exist in the formation of oxides. Firstly, oxides can form
incompletely across the surface of the metal leaving areas unprotected to hrther oxidation.
Oxides of this nature tend to be porous. Secondly, oxides cm form high extemal
compressive stresses so that they flake or break away fiom the metal. This variation is also
unprotective since it leads to exposure of fiesh surfaces. Finally, an oxide can form a
dense, tenacious coating that protects the metal fiom M e r oxidation by limiting the
reaction to ion diffision or conduction. These three variations can be explained through the
use of the Pilling-Bedworth Ratio.
2.2.1 The Pilling-Bedworth Ratio
A method of determinhg whether a metal will form a protective or unprotective
oxide layer is through the use of the Pilling-Bedworth (P-B) Ratio. This ratio compares the
molecular weights and densities of the metal and its respective oxide and is represented by
the equation 2.4:
AOPM P- B Ratio = - AMPO
where A is the rnolecular weight, p is the density, and the subscripts O and M represent the
oxide and metal respectively. This ratio is, in effect, a cornparison of the volumes of the
two materials. If the ratio is less than unity, the oxide has less volume per molecule than
Chapfer 2 Llferaiure Review 5
the parent material and therefore cannot fully cover the metal upon oxidation. This results
in a porous oxide. A P-B Ratio greater than uni@ indicates that there is too much oxide
being fomed across the surface, which results in the development of compressive stresses
and flaking. A ratio of unity allows for protective oxidation. This ratio must be slightly
altered if the stoichiometric ratio between the metal and oxide is not 1: 1. In such a case a
factor 'a' is placed in the denominator. An example of this is in the case of aluminum.
Consider its oxidation reaction.
2 A l + 3 0 , + A 1 , 4 (2.5)
In this case, the P-B Ratio would become:
P - B Ratio = A d ,
2 A ,if PO
Although the P-B Ratio is a good guide that c m be used to predict a metal's
oxidation behavior, it is not completely accurate. Other factors, not taken into account by
the P-B ratio, must be recognized. For instance, similar thermal expansion coefficients of
the metal and oxide will prevent stresses that can arise with temperature changes.
Secondly, oxide plasticity, a high oxide melting temperature, and good adherence between
the metal and oxide will al1 help oxides form a protective layer.
As can be seen from the Table 2.1, a P-B ratio of unity is not essential for
protectiveness, however, values very far away Eom unity are unprotective. Ratios of
particular interest to the present research include aluminum, magnesiun, silicon, and
strontium. Figure 2. I depicts the formation of oxides according to the P-B ratio.
Table 2.1: P-B Ratios of some metals.' Meta1 Al Cu lwg
P-8 Ratio 1.28 1.7 0.84 Pro tective Yes Yes No
Mn 1 2.07
Yes
Si Na Sr 2.04 0.32 O. 69
- Yes No 3
Chapter 2 Litera f ure Review 6
P-B < 1
Tension
P-B = 1 P-B > 1 (Al=l.28)
unifomt iayer
(M0=3.4) compression
Figure 2.1 : Oxide stresses according to Pilling Bedworth Ratios.
2.3 OXIDATION KINETICS The rate at which a metal oxidizes is dependent on several factors, the mechanism
of oxidation, temperature, and composition. In order to study the kinetics of oxidation, the
analysis of weight gain venus time relationship has generally been adopted. The graphical
representation between the weight gain and time can potentially indicate the oxidation
mechanism. Some common oxidation kinetics that have been found experimentally
inchde:*
Parabolic c2 = k, t
Cubic 5) = k, t
Logarithmic 5 =k,alog(a.t+t,)
Inverse Logari thmic 1 - = A - k, log([) 5
where 5 is the oxide thickness as a function of time and, k, a, and A are constants. These
oxide growth equations cover both thin and thick film oxides. However, these equations
are theoretical curves and ofien a metal will display a more complex oxidation behavior,
such as a combination of two or more kinetics, or a change in behavior at different
temperatures. An example of this is tungsten, which has been found experimentally to
Chapter 2 Literature Review 7
exhibit a para-linear behavior, (half parabolichalf linear) at temperatures above 700°C.
Figure 2.2 shows the weight gain versus tirne relationship of severai common kinetics.
O I O 15 20
Time (hm)
Figure 2.2: General classification of weight gain kioetics dunng oxidation.
2.3.1 Thin Films and Scales
The kinetic equations are dependent on the mechanism of oxidation. They are
specific to the various ways in which oxidation takes place for a particular system. It is
generally accepted that the initial monolayer of oxide on metals is formed through the
chemisoption of oxygen ont0 the metal surface. Chemisorption rates tend to be extremely
fast and are therefore dificult to d e t e d e at room temperatures. Measurements that have
been made at low temperatures have indicated a logarithmic relationship. ~ a n d s b e r ~ ~ thus
assumed that the rate controlling step, chemisorption, is therefore logarithmic in nature.
The continued growth of the oxide after the initial monolayer is generally separated into
two categones: thin films and scales.
2.3.1.1 Thin Films
Cabrera and ~ o t t ~ proposed that the growth of thin oxides continues through the
existence of a strong i!ectr;.c field across the oxide. Oxygen ions are formed at the oxide-
Clrapter 2 Literaiure Review 8
gas interface while metal ions are formed at the metal-oxide interface creating an electnc
field which allows for the flow of ions across the oxide. Depending on the conductance of
the oxide, the kinetics can vary.
Oxides can be classified into different types of conductors, n-type semiconductors,
p-type semiconductors, intrinsic semiconducton, ionic conducton, and rnixed ionic
conductors. N-type semiconductors are those which have an excess positive charge,
interstitial cations or anion holes. P-type semiconductors are those that have excess
negative component, cation holes or interstitial anions. Intrinsic serniconductors have
equivalent concentrations of electron defects and interstitial electrons. Ionic conductors
have negligible or very small electron conductivity and conduct either cations or anions,
depending on which has more defects. Mixed ionic also has little electron conductivity but
equivalent cation and anion defects.
The Cabrera-Mott model uses ionic c-ment as the rate determining step with the
assumption that concentration defects are constant within the film. This is probably one of
the areas of discrepancy with this theory since there is fi0 foundation for such an
assumption. The Cabrera-Mott model is parabolic in nature, however, in most cases, thin
films tend to form in a logarithrnic or inverse logarithmic manner. The Cabrera-Mott
model was therefore slightly altered by H a a e and 11schner5 by redefining the ionic current
so that the oxide thickness is dependent on a logarithmic relationship.
As previously mentioned, logarithrnic growth is associated with thin film oxides
and is usually relevant at lower or ambient temperatures only. ~arnmand first deduced this
relationship by observing color changes with the oxidation of various metals in air at low
temperatures. The inverse logarithmic relationship is a much less frequently found
oxidation behavior. Again this behavior is found at low temperatures and only for thin
films in the range of l e s than 100 nm. The inverse logarithmic relationship is often found
to exist at temperatures below or above logarithmic behaviors. An exarnple of this is
tantalun which was found by ~ e m i i l ~ e a ' to exhibit logarithmic kinetics between 50" C and
100°C, inverse logarithmic kinetics between 100°C and 200°C, and a retum to logarithmic
kinetics between 200°C and 300°C.
2.3.1.2 Scales
Scales refer to those oxides that are thicker than 1 O p . 1 and are therefore usually
associated with faster oxidation mechanisms such as linear or parabolic. The easiest form
of oxidation to understand is linear behavior. Oxidation involves the reaction of the metal
ions with the oxygen ions. When there is no bamer separating the two species, oxidation
will occur in a linear manner. Linear oxidation is the most rapid due to the fact that the
reacting ions are readily available. What limits the degree of oxidation in this case is the
surface area exposed during oxidation. A linear relationship with a greater dope will
indicate a larger surface area exposed to oxidation. Metals that tend to oxidize linearly are
those which form oxides that are non-protective md are therefore always exposing new
metal surfaces. Most metals with P-B Ratios below unity and above two oxidize in a linear
behavior due the constant exposure of the metal surface. The resulting oxide tends to ofien
be porous or composed of flakes. Magnesium is an example of a metal that forms a porous
oxide.
In the case of oxides that form in a parabolic manner, the oxidation mechanism
tends to be controlled by diffusion. Initially, oxidation is rapid due to the accessibility of
the metal ions to the oxygen ions. However, as the oxide grows in thickness, the ions must
diffise a greater distance through the oxide to react. The oxidation is therefore effectively
retarded with time. Parabolic kinetics are associated with protective or semi-protective
oxides. In order for oxidation to occur by diffusion, the metal surface m u t be protected by
the oxide. If the oxide is partially protective, or slightly porous, then a thick oxide layer is
required for diaision controlled oxidation.
2.3.1.2.1 Wagner Mechanism
Wagner developed one of the earliest mechanisms for parabolic oxidation of thick
oxides, or scales8. This mechanism is important since it applies to many technical metals
and alloys under seMce conditions. Growth of the oxide in this model occurs through
diffusion dnven by a chernical concentration gradient. The theory is based on the model of
semiconductors where the rate of diffusion is determined by the movement of cations or
Chapter 2 Literature Review 10
anions between vacant lattice sites. In order for directional diffusion to occur, a
concentration gradient, shown in Figure 2.3, is necessary.
Cubic oxidation is not as comrnonly found as parabolic behavior. It generdly
occurs at intermediate temperatures but not for wide temperature ranges. An example of a
metal that undergoes cubic oxidation is nickel at 400°C.
Figure 2.3: Schematic of the diffusion in the Cu-Cu20-O2 systemg.
2.3.2 Dependence on Temperature
It is expected that oxidation rates are dependent on temperature due to the increased
kinetic energy of the atoms associated with higher temperahues. It is now well established
that oxidation kinetic processes follow the Arrhenius equation:
-0 - k = A * e R T (2.12)
where k is the rate constant, A is a constant, Q is the activation energy for oxidation, R is
the gas constant, and T is the temperatce. The dependence of kinetics has been shown to
follow an Arrhenius relationship for many different rnetals. Use of this equation allows for
the determination of the activation energy for the oxidation of various rnetals.
Unfomuiately, due to differing experimental procedures and the extreme precision required
to follow oxidation behavion, large sources of error have led to substantial deviations in
The dependence of parabolic oxidation on temperature is even more obvious. Since
most parabolic relationships are due to a diffusion controlled process, and diffusion is a
thermaily activated process, the effect of temperature is expected. Diffusion can be
expressed mathematically by an equation with similar characteristics to the Arrhenius
equation 2.12:
where D is the diffusion coefficient, Do is a temperature independent pre-exponentid, Qd
is the diffisional activation energy, R is the gas constant, and T is the temperature.
Comparing these two equations, it can be seen that the exponential terms are the same.
Table 7.2 shows some selected metals and how their oxidation behavior changes with
temperature.
Table 2.2: The effect of temperature on the oxidation behavior of various met& (log=logarithimic, lin=linear, parasparabolic, inpinverse, cub=~ubic)~
1 Temp 1 100 200 300 400 500 600 700 800 900 1000 1100 1200 1
log Para para- lin lin 1% Para 1 in lin log log-lin Iin accet
para lin 1 in para para-lin lin-acc
log cub cub para-lin para-lin log cub cub cub cub-lin para para para-lin lin lin accel asym.
1 log inv.log para para-lin lin lin delayed , para para-lin para-lin lin lin Para Para para-lin para-lin para-lin
Fe Ni Cu Zn Al Ge
log log para para Para Para P m Para 1% logcubic para Para Para
log cub para Para Para Para log log para para
log inv.log log para asym lin para para-lin
2.3.3 Alloy Oxidation
It is necessary to consider the effects of alloying elements on the oxidation of a
metal. The concems regarding ailoy .~dditions pertain to their relative reactivity. If the
alloying addition is more reactive &as a higher affinity for the oxidizing speciec) than the
parent nietai, tiizn the oxidation &ha[ wili occur wili foiiow iiit oxidation of hat dioyiny
addition. This c m prove to be both beneficial and detrimental to the oxidation behavior of
the ailoy.
An example of alloy additions being detrimental is the case of aluminum-
magnesiun alloys. Magnesium has a higher affinity for oxygen than aluminum and
therefore will oxidize preferentially before aluminum. The oxidation behavior of
magnesiun is parabolic at low temperatures. However, at temperatures close to 400°C
oxidation becomes linear and non-protective. Aluminum on the other hand is able to
sustain parabolic or protective oxidation to higher ternperatures.
Alloying with an element with a higher afinity for the oxidizing species can also
lower the oxidation rates associated with the pure metal. A cornmon example of this is
stainiess steel, which uses chromium ia retard the oxidation of iron. Chromium combines a
higher affmity for oxygen and a protective oxidation behavior that effectively prevents the
iron from oxidizing at an excessive rate.
A third possibility exists, the reaction or interaction of the two metals to form
heterogeneous oxide layers. Generally separate layers of the compounds of the two metais
or mixtures are formed. An example of this can be seen in the Cu-Ni alloy system that
tends to form two different oxides, Cu20 and NiO.
2.4 OXIDATION OF PURE ALUMINUM This section will deal with the oxidation behavior of pure aluminum at
temperatures ranging from ambient to approximately 1 OOO°C including the kinetics,
mechanisms, and resulting oxides formed.
Chapter 2 Li'terature Review 13
2.4.1 Initial Oxidation
Oxidation of aiuminurn c m be described as being abnormal in that it tends to form
a thin amorphous film at low temperatures. Many studies have reported the presence of
amorphous films ranging 10-30 A in thickness including an early study by ~ i ~ n a r n . " He
f'üïid fhat a plot of tLIc *.vcigDt gain data did not follow a parabolic rclationship but tatlicr a
direct logarithmic relationship for a range of film thicknesses. Dignam developed a
relationship which seems to fully ccver the growth range of the amorphous oxide. This
relationship is s h o w below :
Where X is the film thickness at tirne t, n the concentration of the mobile species, and A
and B are formulae dependent on the conductive nature of the oxide.
Amorphous phases are easily detected by x-ray diffraction analysis. Since
amorphous phases lack any long range order, their diffraction patterns also lack any
noticeable peaks corresponding to a crystalline structure. One of the mechanisms which is
proposed for the initial oxidation of pure aluminurn is that of Bachnch et al.". The model
is applicable for thin films only, a characteristic of the amorphous alumina oxide, and can
be broken down into three stages. The first stage involves the chemisorbtion of oxygen
ont0 the metal surface, followed by the incorporation of the oxygen species below the
surface and the subsequent formation of alumina This mechanisrn was supported recentiy
by Arranz et d l 2 . A more controlled and reliable experimental setup yielded results sirnilar
to ~achrach".
Stucki et al.') studied the initial oxidation of solid and liquid aluminum using
Auger electron and electron energy loss spectroscopy and found that for solid aluminum
only electrons fiom the surface region take part in the bonding to oxygen. It was found that
A1203 begins to form at oxygen exposures in the range of 20 L (1 L = 106 Tom) and the
fmt complete oxide layer ai 100 L followed by a slow incorporation of oxygen underneath
the fmt oxide layer. The h i i n g s were in good agreement with the three stage model
proposed by ~achrach' '.
Chapter 2 Literature Review 14
Initial oxidation of liquid a l h u m differed fiorn solid aluminum in that there
existed an initial phase (-1000 L) before oxide growth took place. Also, much higher
exposure levels (3000 L) were needed to sustain a monolayer of alumina. It was concluded
that in the case of liquid aluminurn, there is no chemisorbed phase of oxygen on the
surface but an uptake of oxygen into the aluminum. The initial reaction is thought to
include oxygen below the surface thereb y excluding ~achrach' s ' rnodel.
The agreement on the mechanism for low temperature (<390°C) oxidation of
aluminum is quite strong. Numerous studies (ranging fÎom 1960 to 1994) al1 confirm the
presence of a thin layer of amorphous dumina which fint forms through chemisorbtion
and continues with completion of a thin layer ranging in thickness fiom 10-30 nm?p8*9 The
thinness of this layer is due to the integrity of the oxide layer and the fact that arnorphous
materials lack any paths of easy diffision such as grain boundaries. This makes ion
diffusion through this material dificult and growth very slow. Although there is sorne
discrepancy as to the kinetics, generally a logarithmic relationship is accepted for low
temperatures. The initial oxidation of liquid aluminum is still questionable. Ody the study
by Stucki et al.13 has focused on the initial oxidation of molten aluminum.
2.4.2 Scale Growth
Oxidation of aluminum at high temperatures diffen fiom that at low temperatures
in that the oxidation is not limited to an arnorphous layer. ~i~narn'~*'~ performed another
study on the oxidation of super purity aluminurn at temperatures in the range of 250°C to
450°C. He found that the oxidation behavior changed at approximately 450°C. At
temperatures below 450°C, the oxidation included the formation of a very thin amorphous
oxide by the logarithmic mechanism proposed by ~achrach". Experiments done at 450°C,
showed that the oxidation kinetics increased after 4 hours. It is believed that this increase
in growth is due to the nucleation and rapid growth of y-A1203 crystallites at the rnetal-
amorphous oxide interface. The growth rate afler approximately 16 hours fmally decreased
off due to what is believed to be crystal coalescence and continued growth consisted of the
formation of a film of these crystallites. The resulting oxidation growth curve could not fit
Chapter 2 L iterature Revie w I S
any one rate law. However, initiai oxidation fits a logarithmic trend whereas the final
oxidation fits a parabolic relationship.
Another early study of the oxidation of aluminum was performed by ~me l t ze r '~
who oxidized super purity aluminum in the 400°C to 600°C temperature range. He found
that the oxidation best followed the parabolic law described by Cabrera and ~ o t t ~ . From
this relationship, the activation energy was detemiined to be 37.2 kcdrnol.
A similar study was performed by Cochran et al." at similar temperatures.
Superpurity aluminum oxidized in a manner which failed to fit any of the previously
identified rate laws. A parabolic relationship seemed to be the closest fit. Weight gains
obtained during these experiments were much lower than other similar studies, such as that
of Smeltzer. Cochran labels this as a surface roughness effect. If the aluminum surface is
too rough, the initial oxide which forms must break due to the large asperities, thereby
exposing fiesh metal surface.
A study at molten temperatures (800-950°C) by Bergsmark et alI8 identified
another variation in the oxidation of aluminum. In this case commerciaily pure aluminurn
was oxidized under various oxygen pressures. Weight gain trends undrr these temperatures
were dso relatively cornplex. The initiai oxide that formed was believed to be y-A1203. In
this case, temperatures were likely too high for the amorphous oxide, found at lower
temperatures, to be stable for any substantial length of time. The samples oxidized at
800°C resulted in what is believed to be a compact oxide layer of microcrystalline y-A1203 4 with the presence of some magnesiurn rich (up to 35%) areas. This would indicate that
some of the oxide was indeed MgO. With tirne, the oxidation slowed down until a point
where the oxide began to tramforin into a-A1203. This was represented by a rapid increase
in the oxidation kinetics and a sudden growth of a porous scde oxide. Samples oxidized at
850-950°C resulted in the detection of only a-A1203 and were characterized by porous
oxide scales which contained some entrapped aiunhum metal. This also showed that
increased oxygen pressures increased the oxidation rates and that wet atmospheres
somehow increased the oxidation behavior of aluminum. An explanation for either
occurrence was not given.
- -
Impey et al.19 also studied oxidation of commercial purity aluminuxn at 750°C
under dry and humid atmospheres. He found that initial oxidation in dry air proceeded
quickly at first, tapering off after a few hours and then accelerating again after
approximately 15 hours. He explained the initiai oxidation was due to the growth of a
crystalline y-A1203 oxide layer, which after time became thick enough to slow down
oxidation. The subsequent increase in oxidation was due to the transformation of y-A.i203
to a-A1203 and the growth of oxide nodules on the surface. Comparatively, the oxidation
in humid air was not nearly as pronounced. Weight gains were approximately seven times
less than those found under dry air. Although oxide nodules were present at the end of the
experiment they were individudly spaced throughout the surface, unlike the dry air
experiments which had a complete coating of nodules. The oxide clusters in this
expenment also contained some entrapped aluminurn metal. An explanation for the
reduced oxidation with humid air was explained by the fact that y-A1203, which has a cubic
spinel structure, tends to be stabilized by the presence of hydroxyl ions. Removal of these
ions by reducing the water vapor pressure results in a much more rapid conversion of y-
Alto3 to a-A1203.
Oxidation of pure alurninum at temperatures above the melting point can be said to
be different fiom oxidation in the solid state. Literature seems to agree that at temperatures
in the range of 750-950°C a transformation Rom y-&O3 to a-Alz03 will take place after
an initiai induction period. A characteristic of this transformation is a decrease in oxide
volume of approximately 24%. This decrease in volume results in the development of
tensile stresses on the oxide and subsequent cracking of the scale. Following failure,
aluminum exudes through the oxide and oxidizes in a breakaway manner.
2.5 ALUMINUM-MAGNESIUM ALLOY OXIDATION
One of the major ailoying additions made to duminum is magnesium. As such its
effect on the oxidation of aluminum is of importance. What makes this alloying element
Chapter 2 Liierature Rewiew 17
considerably more important than others is the fact that it has k e n found to greatly
accelerate the oxidation process of aluminum during heat treating and melting. In order to
understand the effect of magnesium additions in aluminurn, it is necessary to first look at
the physical characteristics of rnagnesium and compare them to alurninum.
2.5.1 Magnesium Oxidation
Magnesium is a metal with a very high vapor pressure such that, at casting
temperatures, alloys with high magnesium contents (>3%) have been found to lose
substantial amounts of magnesium through vaporization. This translates to the fact that
high magnesium containing aluminum alloys cannot be held under vacuum at high
temperatures or too much vaporization occurs. Additionally, magnesium has a higher
affinity for oxygen than alurninum. This can be seen by comparing the heat of formation of
both alumina and magnesia. At a temperature of 298K, the standard heat of formation of
A1203 is 1,113.6 kJ/mol and that of Mg0 1,202.2 k.J/rno~~~. This indicates that if aluminurn
and magnesium are readily available at the rnetal surface, the oxide which would form
should predominantly be MgO.
Since magnesium is the metal that is reacting, it is important to understand the
oxidation behavior of this metal. Previously, it was rnentioned that aluminum oxidized to
form a compact oxide layer at low temperatures. This is supported by the fact that the P-B
ratio for aluminum is only slightly greater than unity (1.28). Magnesium on the other hand
has a P-B ratio less than unity (0.81) and although it is only slightly less than unity, it
indicates that the oxide does not form a complete layer and therefore oxidizes in a non-
protective manner.
The oxidation of pure magnesium was studied by Rosenband et d2'. It was
observed that the initial oxide layer formed at room temperature was protective and was
approximately 70 A in thickness. Upon oxidation at temperatures between 400 and 500°C,
it was found that oxides formed below 475°C remained protective and oxidized in a
parabolic manner. Moreover, at temperatures above 475OC the oxidation changed fiom an
initiai parabolic nature to linear due to the cracking of the initial oxide. The mechanism for
Chupter 2 Liierature Review 18
this linear oxidation is rationalized by considering the stresses needed to fracture the oxide.
Once broken, the fresh metal exposed will oxidize in the same manner as the initial layer.
However, the newly fonned oxide is under a tende stress due to the difference in pressure
fkom the atmosphere and the pressure fiom magnesiurn vapor below the oxide. This stress
is greater than the ultimate compressive stress and therefore continual breaking and
oxidation of fiesh metal results.
2.5.2 Aluminum - Magnesium Oxidation
Reports on the oxidation of magnesiurn containhg alurninum alloys have shown a
sirnilar behavior to that of pure magnesium. The main additional consideration that must
be taken into account is the interaction that occurs between the magnesium and aluminum
metals. For instance, alurninurn and magnesium can oxidize to f o m MgA1204 (spinel)
through the following reactions (Haginoya et aie2?:
Table 2.3: Various Gihb's free energies for the formation of spinel.
Reaction Gibb's Free Energy (kJ) (1 OBK)
However, not d l three reactions are likely to occur. For instance, it can be seen in
Table 2.3 that reaction @ is much less thermodynamically stable than reactions O and Q
and can therefore be discarded when considering the reactions that are taking place during
oxidation. Reaction O is aiso unlikely to occur throughout oxidation due to the fact that it
requires direct contact between alurninum ions, magnesium metallic ions and oxygen. In
most cases, initial oxides prevent this IÏom occunhg. Additionaily, the oxidation of
magnesium to form Mg0 can occur. Availability of excess Mg0 would promote reaction
@ whereas it would compete with the reaction O. Therefore, the spinei reaction that is
most realistic to the oxidation of aluminum-magnesium alloys, is reaction @.
Chapter 2 Litetature Review 19
2.5.2.1 Oxidation in the Solid State
The literature available for solid state oxidation of various aluminum-magnesium
alloys is quite extensive. However, in most studies, the materials have variable
chemistries, are oxidized at different temperatures, under different atmospheres, or tested
using different experirnental procedures.
Early experiments on the oxidation of an Al-3% Mg dloy by ~meltze? showed
that for temperatures below 350°C, the alloy was oxidation resistant and closely followed a
parabolic trend. However at temperatures above 350°C, the oxidation rate changed to a
linear behavior. Two other important observations made were the presence of entrapped
aluminum in the oxide and a black discoloration of the top surface. The oxides formed at
higher temperatures were found to be a combination of spinel and MgO. An aluminurn
difiaction pattern also accompanied these two oxides indicating metal entrapment in the
oxide.
Cochran et alu also looked at the oxidation of a sirnilar AI-Mg alloy, a 5052
wrought alloy with 2.35% Mg and traces of iron and silicon. Experiments were conducted
in the 450-650°C temperature range under dry oxygen and humid air. In this case, the
effect of magnesiurn evaporation was taken into account and it is believed that at higher
temperatures, more evaporation occurs and therefore less weight gain than expected takes
place. Weight gains throughout this study were lower in magnitude than those fiom
Smeltzer's work* by a factor of five to ten. This is accounted for by the smoother starting
surface oxidized in Cochran's work. The logic behind this is that the rougher a surface, the
more likely the resulting oxide layer is to break and cause increased weight gains. An
interesting note in this paper was the effect of wet atmospheres. It was found that oxidation
under humid air was slightly greater than under dry oxygen. In al1 humid experiments
smdl surface blisters developed. It was hypothesized that this blistering helped disrupt the
oxide film and expose fiesh metal.
A continuation of the previous studies was performed by Hine et al? and covered
different oxidizing atmospheres at similar temperatures (440-540°C). Eight different alloy
compositions were studied with magnesium concentrations ranging fiom 0.65 to 7.2 %. Al1
Chapter 2 L iterature Review 20
alloys were foundry alloys and as such contained traces of impurities. Average iron
contents were in the range of 0.2%. In this case the effect of humidity decreased the
oxidation behavior. The explanation for this was accounted for in a similar rnanner for pure
alurninum. Hydroxyl ions stabilize the magnesia layer by increasing the volume ratio.
Again, blisters were found to exist below the surface due to the reaction with water vapor
and it is believed that the evolution of the oxidation behavior fiom parabolic to linear is
related to the breaking of these blisters.
Hine also concluded that the oxidation of Al-Mg alloys is independent of the partial
pressure of oxygen as variable oxygen atmospheres failed to change the weight gain trends
of the alloy. The appearance of a black surface was aiso explained by the presence of
metailic aluminum entrained in the oxide. When the oxide was placed in dilute HCl, some
effervescence was visible indicating the presence of fiee rnetal. With time and the reaction
of the fiee metal, the oxide became lighter in color.
The effect of magnesium content was also outlined. It was found that alloys with
higher magnesium contents oxidized in a more rapid rnanner. The difference increased
with longer oxidation times. This trend is directly related to the fact that magnesium is
preferentially oxidized at these temperatures.
Although the studies by srneitze?, ~ o c h r ; i n ~ , and ~ i n e ~ ~ al1 show that magnesium
is preferentially oxidized, they fail to indicate with any confidence the mechanism of
interaction between the species which are formed, mainly spinel and MgO. The first study
to do so was performed by Ritchie et al?' and focused on the oxidation of a dilute Al-Mg
alloy (1.2% Mg). Firstly, experiments performed at low temperatures agreed with previous
studies in that weight gains were low, characteristic of protective oxidation. He found that
at temperatures above 350°C, Mg0 began to form. Although the fiee energy of spinel
formation fiorn magnesia is hi&, transformation did not occur imrnediately, This was
explained with the assumption that the activation energy was relatively hi&, preventing
the reaction until prolonged heating took place. As longer times were reached, the Mg0
was completely transformed into spinel. One notable problem with this work is that it
assumes that the initial oxide formed on the surface is an amorphous Al2O3 film. Detection
of this layer was not performed and the basis for the oxidation behavior relied on this
initial layer acting as a magnesium ion carrier. It is reasonable to assume that the initiai
layer might actually be an amorphous Mg0 layer. If this is the case, the work by Ritchie is
unsupported.
Field et al.26 continued Ritchie's work with the same assumptions and proposed a
model for oxidation of an Al- 4.2% Mg alloy for the same temperature ranges. It is stated
that at temperatures above approximately 400°C there is a change in growth behavior fiom
protective to nonprotective. This is explained by the nucleation of primary M g 0
crystallites below the amorphous Ai203 overlayer and secondary Mg0 microcrystals
within the Alz03 layer. It was the crystallites that grew and forced the oxide to crack,
exposing fiesh metal and causing a breakaway phenornenon. This paper fails to mention
any formation of spinel at these temperatures. This contradicts the majority of other studies
in this area.
Commody, oxidized Al-Mg alloys form two oxides, Mg0 and spinel. It is agreed
in the literature that Mg0 is the first oxide to form d e r an initial induction period at
temperatures above 400°C. Zayan et al?' showed this to be true as well as the fact that the
outside layer of oxide is always Mg0 and that spinel is formed beneath this top layer by
reaction of aluminum with M g 0 as described by reaction Q in Table 2.3. He proposes that
the mechanism for growth of these oxides is the solid-state diffusion of Mg through the
MgOIspineI oxide constituents. He also found that the mechanism for both a 0.4% and a
2% Mg containing duminun alloy were similar except that at higher magnesium contents,
the diffusion of vapor across voids played a larger role.
zayan2' performed another study that focused on the formation of voids at the
oxide metd interface and concluded that these voids were caused by the outward diffusion
of Mg and resulted in detachment of the oxide. Continued growth of these voids led to
cracking and breakaway oxidation. His model is shown in Figure 2.3.
crack
grain boundaries
Figure 2.4: Mode1 for solid-state breakaway oxidation
2S.L 1.1 Summav of Solid State oxidation
The oxidation of AI-Mg alloys in the solid state resuits in an initial induction period
where the oxide remains protective. Subsequent breaking of this oxide through the
formation of M g 0 crystals at temperatures in excess of 400°C increases the oxidation to a
linear behavior. Spinel is subsequently formed below the M g 0 layer through the reaction
of aluminum and MgO, probably according to reaction (D in Table 2.3. The surface
appearance always has a cauliflower morphology caused by the rupture and oxidation of
surface blisters. Generally wet atmospheres will help reduce initial oxidation through the
stabilization of the present oxide through the formation of hydroxyl groups. However, after
time, oxidation kinetics will increase by aiding the formation of surface oxide blisters.
2.5.2.2 Liquid State Oxidation
Very few oxidatim studies have been performed at temperatures slightly above the
melting temperature and those that do exist tend to contradict one another. Oxidation under
these conditions is expected to be slightly different from solid state oxidation due to the
difference in transport of metallic species in the ailoy. During solid oxidation, ions must
difise through the metal, whereas during liquid oxidation, transport to the surface is much
quicker. Another consideration that must be taken into account is the possibility of melt
disturbance. Since the metal is a liquid, disturbance of the oxide layer is very likely and
rnight have a more profound effect on the underlying metal.
Irnpey et al." perforrned oxidation experiments at 750°C for Al-Mg alioys ranging
in Mg concentration from 0-5 Wto!. As with solid state oxidation, the addition of
magnesium greatly increased the oxidation behavior. Impey believes that the first oxide
present was an amorphous dumina layer that provides some degree of protection.
However, after a short induction period (sometimes as low as five minutes or as long as
five hours) growth of magnssia crystals in the initial arnorphous layer was observed.
Continued growth of the Mg0 crystals resulted in fracture of the oxide and continual
exposure of fiesh liquid metal and rapid formation of oxide. This oxidation process slowed
oniy when the magnesium concentration was too low for Mg0 formation. Similady, as in
solid state oxidation, it was observed that M g 0 formed in a secondary rnanner in the
amorphous film and in a primary manner at the oxidehetal interface. As the oxidation of
Mg0 progressed, the local concentration of Mg in the surrounding area fell below a critical
level so that Mg0 reacted with Al and oxygen to form spinel. The spinel formed in the
shape of large crystals at the oxide/melt interface.
The final oxide layer had a large amount of entrapped metal that was lefi
unoxidized. This was significant in that it showed that magnesium additions can increase
the amount of metal loss through dross formation.
Silva et al.'' also performed sirnila. experiments with slightly different results. In
this case most oxidation expenments lasted approximately two hours. Oxidation
experiments conducted between 650°C and 725°C showed that a protective oxide layer
formed for AI-Mg alloys with less than 3% Mg. Weight gains in this case decreased for the
full two hom. Alloys with Mg contents between 3 and 5% exhibited breakaway oxidation
as described by Impey. An interesthg aspect of Silva's work is the consideration of phase
stability. Figure 2.5 depicts which oxide should be more stable at various Mg contents.
This figure supports the idea that when areas in the melt becorne depleted of magnesium,
the Mg0 will react to form spinel. h contradiction to this is the fact that for d l AI-Mg
alloys tested, only Mg0 was observed. Figure 2.5 indicates that for magnesium contents
below 1.3% spinel should form as a primary oxide. However, spinel is thought to be
dificult to nucleate at the surface and therefore is formed through reaction 8 outlined in
Table 2.3.
Differences in results exid between Impey and Silva. Silva observed through
elemental analysis and x-ray difiction that the initiai oxide forxned was arnorphous Mg0
whereas h p e y believed the initial oxide to be amorphous Alz03.
Chapter 2 L iteraiure Revie w 24
Figure 2.5: Stability diagram for MgO, spinel, and ~ 1 ~ 0 2 ~
The relationship between Mg0 and Mg&O4 was descnbed by Haginoya et It
was found that for alloys containing fiom 3 to 12% magnesium, there was a change in the
oxidation behavior ai approximately 77S°C. Below 77S°C, weight gain cuves exhibited
two waves or humps in the cuve indicating a change in rate controllhg steps, most Iikely
between breakaway and protective oxidation. At temperatures above 775OC, oxidation was
represented by a steeper weight gain cuve with no waves.
Additionally, for temperatures below 1048K, the initial concentration of Mg0
increased briefly and then decreased as magnesium depletion caused the transformation to
spinel. For higher temperatures, Mg0 concentration was low and spinel formation was
found to be high initially and continued to Uicrease with tirne. From these results Haginoya
concluded that for temperatures above 1048K the formation and transformation of Mg0 to
MgN2O4 was very rapid while at lower temperatures it was a much slower process. It was
also observed that both Mg0 and spinel were located at the surface of the alloy; M g 0 in
granuiar f o m and spinel as a layer of fine oxide particles.
Chupter 2 Lit era f ure Review 25
The effects of various atmospheres and pretreatrnents on the oxidation of molten
Al-Mg alloys have also been studied, mainly by Belitskus and Cochran. cochran3' found
that adding beryllium, sodium or dusting boron on the oxidizing surface resulted in an
lengthened induction period before the onset of breakaway oxidation. ~ e l i t s k u s ~ ~
confirmed that boron did prevent breakaway oxidation. In this case, it was found that B203
was incorporated into the M g 0 surface film.
Similarly, atmospheres of CO2 and N2 had the sarne effect. Tests performed under
atmospheres containhg SOz and H2S prevented breakaway o~idation?~ Belitskus
postulated the reason for this being the incorporation of MgS04 into the oxide. It is
believed that S03 is the inhibiting species, however, further studies dealing with off-gas
production within a stabilized atmosphere were recornmended.
Experhents at slightly higher temperatures have been carried out in order to better
understand a process that has found application in the processing of dumina composites.
AI-Mg alloys oxidized in the temperature range of 1100 to 1400K have been found to
exhibit unusually fast oxidation rates. Figure 2.6 depicts how temperature and magnesium
content are related to this phenomenon of rapid oxidation.
A - O
* 1 - A
1100 1200 1300 1400 7 500 Setpoint Temperature (OC)
Figure2.6: Selected data indicating the dependence of the orridation rate on temperature and mapesium content of a Al-Mg-Si alloyM
Chapter 2 Liferature Revie w 26
General agreement exists on the process. Several have s h o w that
rnagnesium is a necessary addition for the rapid oxidation of aluminum. The initiai
breaking of the oxide layer associated with magnesiurn additions allows for the continual
oxidation of aluminum. Generally, it has been found that the top oxide consists of a layer
of M g 0 at th3 oxideigas inkrlca and a iayer of spînei and ennaineci aluminum metal
below this Mg0 layer. It is believed that the continuation of oxidation behavior is due to
the continual feeding of liquid aluminum through the oxide layers. This feeding of
aluminurn is termed ' wicking ' . Studies under these conditions have also shown that the partial pressure of oxygen
is a determining factor in the rate of the reaction. Figure 2.7 depicts the effect of oxygen
partial pressure and temperature on the growth rate of the oxide on an Al-Si-Mg alloy.
O 0.5 1
Oxygen Partial Pressure (atm)
Figure 2.7: Effect of Po2 on the growth rate of a Al-Mg aIIoy3'
Chapter 2 Lr'ferature Review 27
2.6 EFFECT OF SILICON ON OXIDATION Silicon is the most common ailoying element in the aluminum alloy family due to
the increased castability associated with silicon additions. There exists a wide range in the
arnount of silicon additions (045%). This allows for some good tailoring of the desired
dlûy. Few oxidaiion studies exist on binary Ai-Si aiioys, however some information is
available on the effect of silicon on composite formation at high temperattues.
~uermazi~* studied the effects of Mg and Si additions in aluminum and found that
the addition of silicon to pure aluminum failed to change the oxidation behavior. Weight
gains were similar for both alloys. The addition of magnesiurn caused a complete change
in the oxidation behavior, as did the addition of silicon to the Ai-Mg alloy. Silicon
shortened the incubation period for breakaway oxidation. The mechanism of this behavior
is not well understood but it is supported by the fact that application of SiOz particles to the
surface also reduced the incubation period as seen in Figure 2.8.
100 1
Figure 2.8: Effect of silicon additions on the incubation period of a AI-Mg a 1 1 o ~ ~ ~ . ((7/3) = 7% Si / 3% Mg, (OB) = 0% Si 13% Mg)
It has also been observed that Ai-Mg alloys containhg silicon oxidize in a more
consistent maMer than those without s i l i~on?~ Oxide growth is reproducible and steadier
due to the decreased viscosity associated with silicon additions as weil as improved oxygen
supply for oxidation. The decreased viscosity allows for better wicking abilities. Secondiy,
Chapter 2 Literature Review 28
the addition of silicon to aluminum alloys containing magnesium resuits in an increase in
the solubility of oxygen in aluminum, as is s h o w in the Figure 2.9.
Figure 2.9: Solubility of oxygen in aluminum alloys containing magnesiurn with siiicon additiond8
~rnadij' also studied the oxidation of binary Ai-Si alloys and results were similar to
the oxidation behavior of pure aluminum for the temperature tested (740°C). Growth
behavior was charactenzed by two different stages: an initial slow oxidation period
followed by a penod of faster kinetics, indicative of the transformation of y-Ai203 to a-
M203.
2.7 EFFECT OF STRONTIUM ON OXIDATION
Strontium is a common elemental addition used in the aluminum industry to
modifj the rnicrostnicture of silicon containllig aluniinum alloys. Common levels of
strontium additions range between 100 to 300 ppm depending on the amount of silicon
present in the alloy. The effects of strontium are seen mainly in the mechanical properties.
The modified silicon structure has a less acicular shape and is better disûibuted throughout
the alloy leading to improved mechanicd properties.
Unfortunately, there exist some disadvantages to using strontium. For instance, it
has been found that alloys containhg strontium will contain more porosity throughout the
structure. More importantly, strontium is thought to play a role in accelerating the
oxidation of aluminum alloys. It has often been observed that molten metal oxides change
colour after the addition of strontium to the melt, especially if added as pure Sr. Very little
work has been performed on this topic. mad di" did show that for an AI-7%Si alloy, the
addition of strontium resulted in a more rapid initial oxidation (for the f is t 7 hous), but a
more protective long term oxide layer. It was suggested that aluminum-strontium oxides
that formed on the surface were the reason for differing kinetics. Figure 2.10 reveals the
impact that strontium has on preventing the formation of oxide nodules.
It is reasonable to assume that strontium would not form a protective oxide due to
its iow P-B Ratio of 0.69. The presence of a stable oxide layer indicates that strontium is
not the only element involved in the oxidation process.
A) AI-7%Si showing oxide nodules B) AI-7%Si with 250ppm Sr Figure 2.10: SEM pictures of top oxide formed at 740°C for 24 hoursJ0
The oxidation of aluminum alloys is industrially significant in many respects. Solid
state oxidation can lead to surface degradation and liquid state oxidation is one of the
leading causes of metai loss through metal entrapment within the oxide and direct
oxidation. Additionally, rapid oxidation rates at high temperatwes have led to an
Chapter 2 Literature Review 30
application in the processing of ~ a n x i d e ~ ~ ceramic composites through directed melt
oxidation at high temperatures (9504250°C). For these reasons, the study of aluminum
and aluminum alloy oxidation has become a topic of interest to many different industries.
The literature reviewed indicated a good understanding of the oxidation of pure
aluminurn over a large temperature range. Oxidation is charactenzed by a slow growth of a
compact oxide. Transformation of the initial amorphous oxide to various crystal structures
will cause an increase in the oxidation rate. Typically, 450°C is associated with a
transformation to y-alumina, and approxhately 800°C for y- to a-alumina transformation.
The oxidation mechanism of aiuminum-magnesium alloys is completely different
from that of pure aluminum. Magnesium's higher affinity for oxygen leads to preferential
oxidation of magnesium. Resdting oxides formed include M g 0 and MgA1204. n i e ratio of
these two oxides is dependent on the magnesiurn content, temperature and time of
oxidation. Oxidation rates increase with increasing magnesium content and increasing
temperature. The literature available on the oxidation of Al-Mg alloys is scarce at
temperatures used in foundry operations and existing studies oflen indicate contradictory
results.
The effects of both strontium and silicon on the oxidation of aluminum alloys have
not been studied in any particular detail to date. Silicon has been found to promote
consistent composite growth and shorten the incubation time to breakaway oxidation at
high temperatwes. No studies on the effect of silicon at foundry temperatures have been
found to date. The information on strontium is also very scarce. The available literature has
shown strontium to initially increase aluminum oxidation. Prolonged oxidation, however,
results in a compact protective oxide rnost likely due to the formation of Al-Sr-O species.
An important aspect of oxidation studies is the degree to which experimental
procedure plays a role in the results of the individual midies. It is rare to find two
independent studies that have an identicai procedure. Oxidation experiments are very
sensitive to the procedure used and as a result, large discrepancies in results often exist.
PROCEDURE
RAW MATERIALS
Six different ailoys were used for the oxidation expenments in this work. Four of
the alloys were industrial mock alloys that were cast in house From pure starting
materials. The pure aluminum' used to cast the mock alloys was in the form of 1 cm shots
and was 99.9% pure. The silicon' was in the form of 2.5 cm lumps and had a grade of
99.9999% pwity. The magnesiumQ was cut fiom an ingot that was 99.98% pure. The two
industrial alloys tested were cut f?om a 356 ingot supplied by Alcan. A 90% aluminum-
10% strontium master alloy was used for strontium additions. The chemistries of the raw
materials was deterrnined using Ernission Spectroscopy. The results are summarized in
Table 3.1. Microstnictural examination of the alloys was performed to observe the
morphology of the silicon phase and the evidence of strontium.
Johnson Matthey AESXR Group, Brampton, Ontario. O Timminco Metuls, Toronto, Ontario.
Chapter 3 Ekverimental Procedure 31
Table 3.1: Chernical analyais of the raw materiais used.
3.2 MELTING AND CASTING
Pure Aluminum Pellets
As Received
(Weight %)
Casting was performed usir.g a gas fired furnace set to 750°C. The temperature
was monitored using a chromel-alumel K-type thennocouple. A 1 L Sic crucible, which
was first cleaned by rnelting pure alurninum in it, was used for the casting of each dloy.
In the case of the mock alloys, the pure aluminum was melted first and held for
approximately five minutes, followed by the necessary alloy additions. In order to avoid
excessive evaporation, magnesium additions were wrapped in aluminum foi1 before being
submerged with a graphite plunger into the melt. Stimng of the melt was camed out
using a pure alumina rod to reduce contamination. Once the alloy additions had been
made, the melt was lefk for ten minutes, stirred, skimmed, and cast into a graphite mold
with the dimensions s h o w in Figure 3.1.
356 Alloy
As Received
(Weight %)
Balance Balance 0.03 5 6.7 0.065 0.095
c0.003 0.3 5 c0.003 0.008 <0.002 c0.002 <0.002 ~0.005 <0.002 c0.00 1 <0.002 O. 184 <0.003 0.008 <0.003 0.03 1 <0.003 ~0.006
Chapter 3 Ewperimen fai Procedure 32
Metal Fasteners
Figure 3.1: Graphite mold used for casting of alloys.
3.3 SAMPLE PREPARATION In order to have a consistent size and shape, samples were machined fiom the
cylinders with a height of 1.5 cm and a diameter of 3.2 cm. This size was chosen with
respect to the crucible diameter that was used. Al1 faces of the test samples were ground
using 600 grit abrasive silicon carbide grinding paper. The oxidizing surfaces were
ground further to an 800 g i t finish.
3.4 CRUCIBLE FABRICATION Pure alumîna crucibles, with an approximate hal diameter of 3.4 cm and a height
of 2.5 cm, were used for the oxidation experiments. Crucibles were fabricated in house
using a slip casting procedure (Figure 3.2). Crucible molds were fïrst made using plaster
Chapter 3 likverr'mental Procedure 33
of Paris. Plaster and water were mixed with a 70:30 weight percent ratio. A small glass
beaker with an outer diameter of 4.1 cm was used as a pattern for the cnicible shape.
Once solidified, the mold was air dried for no less than 48 hours and cut in hd f in order
to facilitate the extraction of the crucibles.
The slip was prepared using a riixture of A17SG alumina powder', water and a
deflocculant called DARVAN 821A. The ratios, in weight percent, were 70:30:0.3
respectively. This mixture was then bal1 milled for two and a half hours.
The slip was then poured into the mold and left for approximately five minutes
after which the remaining slip was poured out. The crucible was then dried for ten
minutes before being extracted fiom the mold. The crucible was then M e r air dried for
a period of one day.
Crucibles were sintered in a two step process to allow for machining. The first
step included a partial sintenng tirne of one and a half hours at 1200°C. Any anomalies
were easily removed and holes were dnlled into the sides of the cmcibles in order to
accommodate hanging from a wire. Sintering was then continued for two hours at
1500°C. The final inner uiarneter of the crucibles was approximately 3.4 cm.
3.5 OXIDATION EXPERIMENTS Oxidation experiments were conducted using a thermogravirnetric balance
connected to a vertical dense mullite tube k a c e (Figure 3.3). Samples were placed
inside an aiumina crucible and supended fiom a nickel-chromium wire inside the tube
h a c e . AAer making sure the crucible was centered with respect to the rnullite tube, the
fumace was then evacuated to a vacuum of approximately 100 mTorr and filled with
argon several times before a continual 80w of argon was introduced. Gas flow was
monitored with a flowmeter and introduced into an inlet in the TGA f i m e and out
through the vacuum tubing at the bottom of the setup. The h a c e was then heated to the
test temperature, at which t h e air was inwduced into the h a c e and sample weight
readings were comrnenced. The weight change was moritored continuously with respect
* Ahiminum Comparry of America, Pittsburg, Pennsy lvania
cha~ter 3 Ekverimental Procedure 34
8 to time for a period of 45 hours with a Thermo-Gravirnetric Analyzer, or TGA ( C h D-
100). Three test temperatures were chosen for this study, 700,750, 800°C.
(a) Plaster Mold (b)
Figure 3.2: Schematic illustration of the drain-casting process, (a) fi11 mold with slip,
(b) Iiquid extracted through mold leaving compact layer along walls, (c) excess slip
drained, and (d) casting removed after partial drying.
The funiace was equipped with a B-type (Pt-6%Rh vs. Pt-30%Rh) thermocouple,
with an accuracy of f l%, and was connected to a Gultan West 2050 temperature
controller. Six Sic GLOBAR heating elements', arranged in a circular manner, were used
as the heating source. Tlie h a c e temperature was calibrated using another B-type
thermocouple that was placed in the hot zone of the fumace. It was found that for
temperatures between 700°C and 800°C, the temperature in the hot zone was 5OC lower
- --
' CESIWID, Niagra Falls, New York.
Chapter 3 Experimentd Procedure 35
than that indicated by the temperature controller. This was taken into account throughout
the running of oxidation expenments.
In order to produce accurate weight gain data resulting solely from the oxidation
of the alloy, it was necessary to first prrform trial expenments without a sample. This
allowed for the quantification of oxidation occumng from outside sources, mainly from
the wire used to suspend the sample. It was found that the wire wcight gains were
nrpiigibie. Tnis ïactur wab Iùrilirr iiuiiifird Dy r r u h g aireadj u ~ i c l i ~ d hircs ds iiianq
times as possible.
Micro- Balance
1 Gas
V
Valves
Vacuum Pump -
Figure 3.3: Experimental set-up to measure the oxidation weight gains.
Cha~ter 3 Ex~er im enta1 Procedure 36
3.6 SAMPLE ANALYSIS 3.6.1 Spark Emission Spectroscopy
Chemical analysis of the alloys was quantified using an emission spectroscope
Spectrovac 1Oûû equipped with an MC20 data processing system. Analysis were made at
three areas of the machined casting and an average was taken as the ovenll chemistry.
Figiire 3.4 shows the location of analyses.
Analysis # 7
Analysis # 2
Figure 3.1: Location of samples taken from each alloy Tor chemical analysis
3.6.2 Optical Microscopy
Optical microscopy was used to examine sample microstructures, both before and
after oxidation experiments. Most samplrs were cut using a band saw with the exception
of sarnples with high oxide contents. mainly the AI-3%Mg samples. In the case of highly
oxidized samples, cross sections were cut using a Beuhler ISOMET Low Speed Saw
equipped with rt diamond wafering blade. Beuhler low viscosity, slow setting (24 hour)
epoxy was used for samples that required mounting. The surface was then g o n d down
to an 800 Gnt finish using S i c grinding papers. Polishing stages included: Firstly, 1 um
diamond paste suspended in a water bastid diamond rxtcnder, followed by 0.5 ym
colloidal silica. Ultnsonic clraning was perfomed out before and after each polishing
stage.
Chapter 3 Experimentaf Procedure 37
3.6.3 Scanning Electron Microscopy
Both oxide surface morphologies and polished cross-sections were examined
using the JEOL JSM 840A scanning electron microscope. Qualitative chernical analysis
of surface species was also performed for several of the test samples using Energy
Dispersive Spectroscopy (EDS). Chemical analyses of cross-sections were performed
using Wavelengh Dispersive Spectroscopy (WDS) on an Electron Probe Microscope that
was capable of quantimg the presence of light elements such as oxygen. Due to the
non-conductive nature of the majority of samples, surfaces were plasma coated with
carbon prior to examination with either electron microscope.
3.6.4 X-Ray Diffraction @RD)
X-Ray Diffraction was used to determine the top oxide layers fonned throughout
the oxidation experiments. Two methods were used to examine the samples. Firstly,
surfaces were analyzed as is since the oxides were too thin to remove. Secondly, highly
oxidized samples were leached of the majority of remaining metal by submerging the
sample into a 5% HF solution for twenty-four hours. The oxide remaining was put into a
shatterbox for approximately 15 seconds. tifter the remaining small flattened metal
pieces were separated mancaily, the resulting powder was analyzed by x-ray diffraction.
X-ray diffraction peaks obtained from the as-oxidized samples had high
background noise and slightly shified peaks. The diffraction procedure utilized a scan rate
of 0.2 Ols with filtered Cu-Ka radiation at an accelerating voltage of 40 kV and a beam
current of 20 mA. The diffractometer used was a Philips M D 1700.
The results have been separated into sections pertaining to individual alloys. Ail
results refer ta samples oxidized over a perbd of 45 h o m unless otherwise indicated.
This section will present the resdts directly observed. In-depth details and explanations
will be left to the discussion section
4.1 ALLOY COMPOSITIONS Emission spectroscopy was conducted in order to determine the chemistry of the
alloys used for the oxidation experiments. The results are show below in Table 4.1.
Table 4.1: Composition (in weight %) of the vanous aluminum aiioys used
Synthetic
Allo ys
Commercial
AUoys
Pure Al 99.9 0.03 0.06 c0.003 <0.002
413 87.6 11.1 0.08 <o,oo3 <O.OOS
413 w/ Sr 87.2 11.6 0.07 <0.003 0.03 5
5000 series 95.9 .O4 .O6 4.1 <0.002
356 93.2 6.7 0.095 0.35 c0.002
356 wl Sr 93 .O 6.9 0.09 0.33 O ,023
Chapter 4 Results 39
4.2 PURE ALUMINUM The oxidation experiments for pure aluminurn showed very little weight gain
throughout the full 48 hours tested. Figure 4.1 summarizes the weight gains at 700 and
800°C. No rate law could be fitted to the oxidation curves due to the slightly wavy nature
of the curves. There was little noticeable difference in the weight gains at the two
temperatures tested. No one behavior f i t either curves.
Time (hrs)
Figure 4.1: Weight gain curves for pure aluminum at 700 and 800°C.
Upon rnicroscopic examination of the surface, it was found that a coherent oxide
layer was present on d l samples at al1 ternperatures. Figure 4.2 is an example of the oxide
observed. No surface porosity could be seen on any of the oxides formed. The oxide is
thin enough that the pre-testing grinding lines are still visible.
Due to the rough nature of the solid sanples examined, x-ray diffraction patterns
obtained fiom the pure aluminum samples had large amounts of background noise. All
pure aiuminum samples anaiyzed showed diffraction patterns for metallic alufninum only
with no detectable presence of oxidation products, as indicated in Figure 4.3.
Chapter 4 Results JO
Figure 4.2: Surface oxide of pure aluminum sample oxidized at 750°C.
10 20 30 40 50 60 70 80 90 100 Angle c20)
Figure: 43: X-ray dnraction pattern of a pure aluminum sample surface oxidized at 800°C.
Chapter 4 Resuits 41
4.3 SYNTHETIC 413 ALLOY In a similar marner to pure aluminum, the 41 3 alloy expenenced little in the way
of weight gain, with the exception of the samples tested at 800°C. Weight gain curves are
shown in Figure 4.4, which are uniquely shaped. A transition seems to be occumng
between 700 and 800°C as the shape of the curve changes fiom a basic parabolic curve,
to a wavy curve to a para-linear curve.
0.006
O I O 20 30 40 50
Time (hrs)
Figure 4.4: Weight gain curves for synthetic 413 alloy at 700,750, 800°C.
Surfaces examined revealed the presence of a coherent oxide layer, similar to that
obsewed on the pure aluminum samples. However, samples oxidized at 800°C contained
oxide clusters, approximately 10 to 20 p in diameter, distributed randornly across the
top oxidizing surfaces. Figure 4.5 shows an example of the coherent oxide layer, which
again shows the pre testing grinding lines, and an oxide nodule.
a) Coherent surface oxide (750°C) b) Oxide cluster (800°C) Figure 4.5: Oxide surfaces of the synthetic 413 alloy.
X-ray diffraction patterns for the synthetic 413 alloy also failed to reveal the
presence of any oxide. Patterns obtained revealed the presence of metallic aluminurn and
silicon only. Figure 4.6 shows a
sarnples.
typical x-ray diffraction
A
l 1 ' 1 1
O .-+uLI- I O 20 30 40 50 60
Angle ('20)
pattern obtained for these
A A: Al
Figure 4.6: X-ray diffraction pattern of 413 mock aiioy oxidized at 750°C for 45 hours.
Chapter 4 Results 43
4.4 SYNTHETIC 413 ALLOY CONTAINING Sr The total weight gains associated with the synthetic 4 1 3 alloy containlng 3 50 ppm
Sr were very similar to those observed for the synthetic 413 alloy except for samples
oxidized at 800°C. Lower weight gains were observed for Sr containing samples at
800°C. Weight gains were again very low overall, indicating a protective oxide layer.
O 10 20 30 40 50
Time (hrs)
Figure 4.7: Weight gain curves for the synthetic 413 alloy containing 350 ppm 3r at 700,75O,80O0C.
Optical microscopy was performed on pre-tested sample cross sections of both
synthetic 413 alloys in order to determine if the Sr had modified the silicon structure and
was present throughout the alloy. Figure 4.8 shows the effect of Sr on the silicon
morphology.
Chapter 4 Re s u lts 44
a) Unmodifïed Figure 4.8: Micrographs showing the effect of Sr as a modifier in the 413 mock
alloy.
The surface of al1 samples oxidized wrre similar in appeanncr. Two dii'frrent
colours were observed throughout the surface of the samples: bnght. shiny silver spots
scattered over a dark, shiny surface. Observations at higher magnitïcations under the
scanning electron microscope failrd to reveal any differences in oxide colour or
chemistry. Microscopie examination of the surfaces revealed a cohererit oxide layer ai al1
temperatures tested. Samples oxidized at 8ûû°C failed to reveal any anomalies such as
oxide clusters (Figure 4.9).
Chemical analysis of the oxide layer using an EDS detector was conducted at a
low accelerating voltage (SkeV) to lower the depth of penetration of the electron beam
and to determine if Sr was incorpontrd in the top oxide layer. Estimated depth of
penetntion of the electron beam was approximately 3 ym. Results showed that Sr was
present in small amounts. The elements Al, Si, and O were also detected. Chemical
analysis results from EDS are &en in Appendix 1.
X-ray diffraction patterns revealed the presence of metallic aluminum and silicon,
as shown in Figure 4.10. Some srnall additional low angle peaks were also present. Due
to the existence of only three peaks, their identification however was inconclusive.
Figure 4.9: Microscopy of the oxide surface of the synthetic 413 alloy containing 350 ppm Sr oxidized at 8ûû°C.
10 20 30 40 50 60 70 80 90 100 Angle ('20)
Figure 4.10: Kray diffraction pattern of synthetic 413 alloy containing 350 ppm Sr oxidized at 750°C for 45 hours.
Chapter 4 Resulls 46
4.5 COMMERCIAL 356 ALLOY The oxidation behavior of the 356 alloy indicated much greater weight gains than
both the 413 alloys and pure duminum. Weight gains increased steadily throughout the
45 hour penod for samples at al1 temperatures, as shown in Figure 4.1 1.
20 30
Time (hrs) Figure 4.11: Weight gain curves of the commercial 356 alloy at three different
ternperatures.
All samples tested resulted in substantial visible oxidation. Surfaces were light
gray in color and very cnunpled. The degree of oxidation changed visibly with increasing
temperature. Samples oxidized at higher ternperatures (Le. 800°C) tended to have a more
cnimpled oxide layer and more sub-surface porosity (Figure 4.12). Al1 samples revealed
the same rnicroscopic surface layer which had a 'cauliflower-like' morphology (Figure
4.13).
Chavter 4 Resu lts 47
a) 700°C b) 800°C Figure 4.12: Low magnillcation of top oxide surfaces and of the commercial 356
alloy.
Figure 4.13: High magnification micrograph of the surface oxide on commercial 356 alloy samples.
Again al1 samples had the same exterior color and it was expected that this would
lead to the presence of a similar exterior oxide. Due to the extreme roughness of the top
oxidizing surface, the bottom of each sample was used to obtain x-ray patterns. Al1
samples resulted in spectra of alurninum, silicon, and spinel. An example of an n-ray
pattern obtained is s h o w in Figure 4.14.
Chapter 4 Resulis 48
! I WC' ii
1 I 1 B N C
1 1 1 IBIC Ci; B =CA i I i : i 1 !
i..-J.i\uiiud 1 ! i V I 1 l I O I O 20 30 40 50 60 70 80 90 IO0
Angle ('20) Figure 414: X-ray diffraction pattern of the commercial 356 alloy oxidized at
750°Cm
4.6 COMMERCIAL 356 ALLOY CONTAINING Sr The 356 alloy containing strontium displayed completely different oxidation
characteristics from the 356 alloy without strontium. Weight gains were in the order of
ten times less with the addition of strontium. Oxides failed to have the du11 gray
appearance of the plain 356 alloy, but rather a shinier, colored oxide. Weight gains at al1
temperatures were very similar to each other with a slightiy greater weight gain with
increased temperature, as shown in Figure 4.15. Total weight gains were very similar to
those obtained for the synthetic 413 alloys. The oxidation curves best fit a logarithmic
relationship.
Chapter J Results 49
Time (hrs) Figure 4.15: Weight gain cuwes For the commercial 356 tUloy containing 250 ppm
Sr at three different temperatures.
In order to contirm the presence of strontium, the microstructures of the two 356
alloys were compared. It was observed that the silicoii rnorphology was completely
modi fied, as shown in Figure 4.16.
a) Unmodified b) Modified Figure 4.16: Micrographs showing the modification of the silicon phase in a
commercial 356 alloy due to the presence of 250 ppm Sr.
Cha~ter 4 Results 50
Examination of the top oxide showed very little evidence of heavy oxidation. Al1
surfaces exhibited a dense, coherent surface oxide with the presence of sporadic and
randomly distributed oxide 'extrusions' (Figure 1.17). An increasing number of these
localized areas of oxidation were found wi th increasing tempenture.
Hi gh rnagni t ' k t ion observations revealed that the dense coherent oxide appeared
to be slightly thicker and rougher at higher temprntures, as shown in Figure 4.18. The
locaiized oxide eruptions were characterized by having heaviiy onidizeci surfaces (Figure
4-19), and in some cases, small crystals of magnesia. scattered over the surrounding
surface OP the eniption. The heavily oxidized surfaces of the eruptions greatly resembled
those found throughout the whole oxidizinp surface of the commercial 356 alloy
containing no Sr.
Coherent Oiuide \
Localized Eruption /
Figure 4.17: Low magnification of commercial 356 alloy containing 250 ppm Sr oxidized at 750°C.
a) 700°C b) 800°C Figure 4.18: Coherent surface oxides formed on commercial 356 alloy containing
250 ppm Sr.
Chapter 4 Results 51
Figure 4.19: Localized oxide eruption (7ûû°C).
Examination of the cohrrent top oxidr through EDS showed the presence of
magnesium, silicon, strontium, and oxygen. The intensity of the magnesium peak was the
hiphest, with the oxygen peak being the second in intrnsity. Small peaks of silicon and
strontium were also observed, however, no aluminurn peak was present (Appendix 0.
X-ny diffraction results showed the presrncr of metallic aluminum and silicon.
Some small additionai peaks were unidentifiable (Figure 4.10). The presence of MgA120J
was not Found in any of the samples.
10 20 30 40 50 60 70 80 90 100 Angle c20)
Figure 4.20: X-ray diffraction pattern of commercial 356 alloy containing 250 ppm Sr.
Chapter 4 Results 52
4.7 SYNTHETIC 5000 SERIES ALLOY The oxidation of this alloy was completely difKerent fiom al1 other alloys in that
the weight gains were significantly larger for this alloy. It was observed that after an
initiai penod of little weight gain, the metal began to oxidize in a very rapid marner. The
continuation of this oxidation was very dependent on the test temperature. At 700°C, no
npid midation pc;iod \vas c-dciit. At 750°C, saiipk wcight gains werz csireniziy
complicated. An initial slow growth period was followed by a rapid increase in oxidation
rate. Continuation of this rapid oxidation was characterized by a penod of varied growth
rate that eventudly diminished to a negligible weight gain rate. At 800°C, weight gains
were sirnilar to those observed at 750°C with a few variations: Firstly, the initial penod
pnor to rapid oxidation was, on average, shorter than that observed for the 750°C
samples. Secondly, the rapid oxidation period, was less erratic in its continuation. Finally,
the arresting of the rapid oxidation penod was more abrupt at 800°C.
Al1 samples tested exhibited the sarne rough, black surface. It was evident fiom
the appearance of the samples that at higher temperatures, more of the metal was
consumed through oxidation. The final shape of the samples was increasingly abnormal
with increasing temperature. The samples tested at 700°C retained a relatively Bat
surface, while samples tested at 750 and 800°C were extremely rough.
Due to the extreme roughness of the samples, calculating the weight gain per unit
area is inaccurate. Therefore, witl the assumption that magnesiurn is preferentially
oxidized, graphs of specific weight gain venus time were constnicted (Figure 4.21). This
method quantifies the arnount of Mg that has oxidized to form spinel (the highest oxygen
containing species).
In order to observe the oxidation behavior of the 700°C experiments, Figure 4.22
has been plotted on a lower specific weight scale. It is evident that the 700°C samples
undergo a characteristic oxidation behavior, although on a much smaller scale.
Cha~ter 4 Results 53
O 5 10 15 20 25 30 35 40 45
Time (hrs)
Figure -1.21: Specific d g h t gains for the synthetic 5000 series alloy at three different temperatures.
0.0 O 5 10 15 20 25 30 35 40 45
Time (hrs) Figure 4.22: Specific weight gain of the synthetic 500û series ailoy at 700°C.
Chapter 4 Results 54
Surface Analysis
Al1 samples reveaied similar surface oxide morphologies. Their shape was
characteristic of a 'cauliflower', as seen in the commercial 356 alloy. Al1 surfaces were
noticeably porous, with increasing amounts of porosity at higher temperatures. Figure
4.23 shows examples of the rough oxide, at both low and hi& magnifications.
Due to the extreme oxidation of the 5000 series alloy, intempted nxidation
expenments were conducted at 800°C in order to observe the initiation of rapid oxidation.
Samples were oxidized to points immediately before and after breakaway oxidation
occurred as s h o w in Figure 4.24.
Observation at al1 interrupts reveaied a light gray coating over the majority of the
top oxidizing surface. Microscopie examination revealed the presence of small spikes of
rnetai protruding from the surfaces of the samples (Figure 4.25). The sample that was
oxidized past the breakaway point of oxidation showed that initiation of the breakaway
oxidation occurred dong the sides of the sample and not dong the top surface.
Cross-sectional Anabsis
Cross sections were cut fiom the samples in order to observe the penetration of
oxidation throughout the sample. The oxidation behavior of the sarnples oxidized at
700°C was completely different from that of the samples oxidized at 750 and 800°C.
Oxidation of the 700°C sarnples was restncted to the surface of the simple, whereas large
oxide clusters were observed deep beneath the oxidizing surfaces of the 750 and 800°C
sarnples. Clusters were fmt observed in samples that had entered the area of rapid
oxidation. Qualitative observations of the samples revealed that the number of clusten
grew in size and number with tirne. All oxide clusters were located circumferentially
around large pores in the structure.
In addition to the oxide clusten, oxidation was present in the form of a thick
oxide surface layer (for 750 and 800°C samples). These layers varied greatly in thickness,
frorn 5 to 100 p.
Chapter 4 Results 55
Figure
a) low rnagnification
b) high magnification
4.23: Micrographs showing cauliflower-like morphology of surfac the synthetic 5 0 0 series alloy at 750°C.
Chapter J Results 56
O 1 2 3 4 5 6 7
Time (hrs)
Figure 4.24: Weight gain graph indicating location of interrupt experiments conducted at 8ûû°C.
Figure 4.25: Spikes penetrating the top oxide surface. (800°C, 3 hours)
Chapter 4 Results 57
Chem ical Analysis
Oxide clusters were examined using WDS and EDS detectors in order to
characterize the elements present throughout the structure. Clusters examined after 45
hours of oxidation were composed of only two phases throughout the whole structure,
spinel and aluminum (Figure 4.26). Oxide clusters rxamined after 22 hours of oxidation
were composed of two different sections: an inner circular area that was greatly oxidized
and a surrounding arra irss drnseiy oxidired (Figure 4.27aj. Tir burder briwecii ilicse
two sections is quite visible in al1 cases. The geometry, however, varied slightly with al1
c!usters having at least some Form of rounded shape.
The inner most area of the cluster was composed of two sepante phases:
MgA1204 and metallic aluminum. The second rcgion of the cluster was ülso composed of
two phases: M g 0 and metallic aluminum. In both cases, the metal present was
completeiy devoid of rnagnesium. Figure 4.17 indicates the elemental distribution
throughout one such cluster. (EDS spectrums are seen in Appendix I)
Figure 4.26: Micrograph showing the spinel and aluminum phases present in an oxide cluster oxidized for 45 hours.
Chapter J Results 58
Figure 427a: Secondary electron image of compleie cluster
b) Al 4 Mg 4 O Figure 4.27: Elemental maps showing the distribution of Al, Mg, and O throughout
an oxide cluster a) secondary electron image of entire cluster, b-d) elemental maps of a portion of cluster (750°C, 22 hrs oxidation).
c h a ~ f e r # Results 59
X-ray Diffraction
Diffraction patterns were obtained for the 5000 alloy using two different methods.
Firstly, samples were analysed in the as oxidized condition. Only samples at 700°C
resulted in a readable x-ray difiaction pattern. Results shown in Figure 4.29a indicate
that the surface oxide present on the 700°C samples was M g 0 with some aluminum
present. Secondly, samples were leached of their rernaining metallic componenis by
submerging portions of the sarnple into a 10% HCI solution and puiverizhg the product
using a shatterbox. The resulting powder was cleaned and analyzed. Results from this
method (shown in Figure 4.28b) indicate the presence of MgA1204 and aluminwn. The
presence of Mg0 was not found using this method.
d o 20 30 40 50 60 70 80 90 100 Angle ('20)
a) Surface scan
AIC
1
Powder A: Al C:MgAI2O4
? O 20 30 40 50 60 70 80 90 100 Angle ('20)
b) Milled powder Figure 4.28: X-ray diffraction patterns for the synthetic 5000 series alloy.
5. DISCUSSION
5.1 OXIDATION OF PURE ALUMINUM
The pure aiuminum samples exhibited minima1 oxidation as is evidenced by their
weight gain curves shown in Figure 4.1. Oxides formed are considered protective due to
the fact that after prolonged exposures of 45 hours, weight gains did not significantiy
increase. Weight gains in the range of 0.003 g/cm2, when translated to oxide thickness,
result in a top oxide layer of approximately 4 pm (asssuming oxidation of the top of the
sarnple only). This thickness is an over-estimate due to a miniscus effect. The sample
surface was slightly curved and therefore, more surface area was exposed to oxidation. In
addition, it is impossible to lirnit the oxidation to the top surface of the sample.
The fact that x-ray diffraction was unsuccessful in detecting an oxide layer
strengthens the point that the layer is very thin. It is assumed that y - N z 0 3 is the oxide that
forms. Literature suggests that alumina layers that undergo a transition fiom y to a tend
to show signs of oxide racki in^'^*^^. This cracking is due to a decrease in respective oxide
volume and is associated with the formation of oxide nodules at the cracks. Oxide
nodules were not observed on any of the pure altmiinum samples. Additionaily, a
transformation fiom y to a translates to change in the weight gain curves. Oxides
composed of y-alumina generally exhibit a parabolic weight gain relationship whereas a-
durnina exhibits a linear relationship. No distinctive change in the weight gain curve was
observed throughout the oxidation of the pure alunhum samples and therefore it is
believed that a protective y-A1203 layer exists on al1 samples tested.
cha~ter 5 Discussion 62
This conclusion contradicts slightiy some of the current literature that states that
the y to a transition temperatures is in the range of 8 0 0 ~ ~ ~ ~ ~ ' ~ . The reason that no
transition is observed might be due to the presence of slight amounts of water vapor. It
has been observed in previous stuclies that hydroxyl ions are y-A.i203 stabilizers.
Therefore, the presence of water vapor would delay the transition fiom y to a dumina.
5.2 EFFECTS OF SILICON
The synthetic 41 3 alloy was tested in order to determine the eflects of silicon on
the oxidation behavior of duminum. It was found that silicon increases the amount of
oxidation at 800°C. Synthetic 413 alloy samples oxidized at 800°C showed much larger
weight gains than the pure alurninum samples (Figure 5.1). Weight gain differences at the
two lower temperatures of 700 and 750°C were minimal and can be accounted for by the
resolution limit of the microbalance.
O I O 20 30 40 50
Time (hrs)
Figure 5.1: Cornparison of weight gains of pure alurninum and 413 aLioy at 800°C.
In addition to the sirnilarity in weight gains, both the pure alufninum and the 413
alloy displayed the same top oxide morphology, with the exception of the 413 alioy
Chapter 5 Discussion 63
oxidized at 800°C. Cohesive, protective oxides similar to those found on the pure
aluminum samples were observed on the 4 13 sarnples. These samples, oxidized at 800°C,
were quite different to others due to the presence of oxide nodules, as shown in Figure
4.5b. The nodules were very similar to those observed by ~ r n a d i ~ ' (shown in Figure
2.10a) who tested a Al-7%Si alloy at 740°C. These nodules are most likeiy composed of
a-dumina with entrained molten aluminum. The reason that nodules fonned on the 413
dioy and not on the pure aiuminum may be due ro the ciecreased viscosity associated with
the addition of silicon. If a small crack forms in the top oxide layer, it is easier for the
metal to feed through the cracks and to continue to oxidize.
Cracking of the oxide layer is attributed to localized transformation of y to a-
alumina. As stated previously, a-alumina oxidizes in a linear manner through the
formation of nodules. As a-alumina begins to form and the initiai y-alumina layer cracks,
molten metal flows through the crack and fiesh metal is exposed to the oxidizing fiont.
The resulting oxide formed fiom the freshly exposed metal is no longer flat and lacks the
strength and integrity of the initial oxide. This allows for continued cracking and
Iocalized growth normal to the sample surface. The result is an oxide nodule exhibiting
Iinear oxidation kinetics. The growth curve that was observed for the synthetic 413 alloy
at 800°C is not linear. This c m be explained by the sporadic appearance of the oxide
nodules. The average size of these nodules was approximately 10 pn, and they were
scattered randomly across the oxide surface. The rounded shape of the oxidation curve is
a result of the combined effect of a slow thickening of the y-alutnina layer that covers the
majority of the sample
a-dumina nodules.
5.3 EFFECTS ALLOY
suface and the continual growth of the small sporadically placed
OF STRONTIUM ON SYNTHETIC 413
The synthetic 413 ailoy containhg 350 ppm strontium was tested in order to
examine the effects of strontium on the oxidation of an Al-Si alloy. It was found that
samples tested had similar weight gains to pure duminum. Oxide surfaces were darker in
colour with a shinier surface. Microscopie analysis reveaied a protective oxide layer
Chavter 5 Discussion 64
throughout the samples. Samples tested at 800°C failed to f o m any oxide nodules on the
surface. This was supported by comparing the weight gains of pure aluminum, and the
two synthetic 413 alloys at 800°C as shown in Figure 5.2. It can be seen that the 413
dloy without strontium has the highest weight gain.
Figure 5.2 has been tmcated to show the initial oxidation period. Observations
made by ~rnadi~ ' indicate that strontium additions increase the initial oxidation rate, but
decrease the total oxidation in the long tem. As is evident in Figure 5.2, this not mie in
this study. Oxidation rates were slower throlighout the entire oxidation t h e for the
strontium containing alloy. One of the reasons that this might occur is that strontium
plays some role in stabilizing the dumina layer that forms. Another hypothesis is that
strontium forms an oxide with aluminum or silicon, as suggested by Emadi. Either of
these mechanisms might prevent or delay the transformation of y to a-alumina and
therefore prevent the cracking of the oxide and subsequent nodule formation.
O 1 2 3 4 5
Time (hrs)
Figure 5.2: Cornparison of the initial oxidation behavior of pure aluminum, 413, and 413 containing strontium.
Chapter 5 Discussion 65
5.3.1 X-Ray Diffraction Analysis
An interesting point can be made by examining the x-ray difnaction results.
Figure 5.3 compares the low angle x-ray difiction patterns of the two synthetic 413
alloys. The diffraction patterns have been tnuicated to better distinguish the low htensity,
low angle peaks obtained. Low angle peaks are characteristic of a large ce11 stucture. It
cm be seen that several small, unidentified peaks are present in the strontium containing
aiioy. Aithou& no perfect match was made with any Al-Si-Sr-O compounds, it is
reasonable to assume that these peaks are due to some fom of strontium containing
oxide. Emadi, in his snidy reported x-ray diffraction peaks fiom SrA1204 and Sr&07.
10 - 15 20 25 30 35 40 45 50
413 containing 350 ppm Sr 4
4- 1
L - 1
r
+ 1, I I
I
10 15 20 25 30 35 40 45 50
Angle ('20) Figure 5.3: Cornparison of x-ray diffraction peaks of 413 and 413 containing
strontium alloys.
The peaks that are present in Figure 5.3 corne close to matching some oxides,
such as Sr&07, SrSi03, and SrA1204, as well as a strontium intemetallic, AlzSirSr.
Chapter 5 Discussion 66
However, several peaks are missing fiom each pattern It is difficuit to match any of the
oxide dif ic t ion patterns due to the problems associated with the examination of non-
planar sample surfaces. Rough sarnples ofien result in shifted peaks and preferred
orientation. This can lead to certain peaks having abnomally high intensities compared to
others.
5.3.2 Microanaysis
In order to conf~rrn the presence of strontium in the top oxide region, elemental
scans of the top oxide surface were made by EDS. It was found that strontium was
present in the top 4 pm of the surface, as well as aluminum, silicon, and oxygen. The
EDS analysis shows very strong alurninum and oxygen peaks. This confïrms that the
majority of the top oxide layer is alumina. It is possible that the strontium detected is due
to the presence of the aforementioned silicide, A12Si2Sr. This, however, fails to explain
the effect of strontium on the oxidation behavior. Furthemore, formation of the
intermetallic should not affect the oxidation behavior unless it acts as a nucleating agent
for oxide formation.
Another hypothesis is that strontium forms an oxide with alurninum and silicon
fiom the initial alumina layer formed on the surface of the sample at room temperature.
Strontium is known to fade with tirne through oxidation. Therefore, perhaps a strontium
rich region forms below the initial oxide and reacts to form strontium containing oxide.
This secondary oxide rnight not only be protective, but might also prevent a-alumina
fiom forming. This hypothesis would explain the presence of intense alurninum and
oxygen peaks, which indicate the presence of alumina, observed during EDS analysis.
Strontium plays some role in preventing the formation of oxide nodules. The
exact mechanism still remains uncertain. Previous hypotheses that strontium increases the
oxidation rate of Al-Si alloys are unsupported in this study.
5.4 OXIDATION OF SYNTEZETIC 5000 SERIES ALLOY
Oxidation expenments were performed on this alloy in order to observe the role
that magnesium plays on the oxidation behavior of alUrninum. It was observed that
Chapter 5 Discussion 67
weight gains were dependent on the temperature of oxidation. Samples oxidized at 700°C
displayed weight gains that were low compared to samples oxidized at 750 and 800°C.
Ail oxidized samples were black in color. This is believed to be due to the presence of
entrained aluminum near the surface of the samples, a contention that is supported by
previous studies performed on Al-Mg alloysuv24. The high weight gains observed are
directly associated with the presence of magnesim. It was found, through x-ray
di f icdon mdysis, thet only rnapesirur. condnhg oxides xc re prescrit, bath on, anJ
below the surface, in any of the samples.
Oxidation of the 5000 series will be charactenzed by breaking down the oxidation
behavior into separate stages: an initiai slow oxidation stage, followed by a stage of rapid
oxidation and subsequent continuation or discontinuation of rapid oxidation.
5.4.1 Initial Oxidation
Initiai oxidation is defined as the time required to Uiitiate breakaway oxidation, or
the induction penod. The weight gains associated with the induction period are quite low
compared to other stages of the oxidation process. By o b s e ~ n g the specific weight gain
curves shown in Figure 4.21, it can be seen that as the temperature increases, the
induction penod decreases. Table 5.1 gives the average induction period for the three
temperatures tested.
Table 5.1 : Average induction periods.
II Temperature 700°C 750°C 8OO0C 11
From the above table it is obvious that there is a significant difference in the
behavior between the samples oxidizeà at 700°C and those oxidized at 750 and 800°C.
Samples oxidized at 700°C never entered the second stage of oxidation. Figure 5.4
depicts a typical weight gain curve of samples oxidized at 700°C. The cume is broken
down into three separate stages: an initial parabolic curve followed by two linear
portions.
Induction Period (hou rs) ..'
45 10 2 1
Chapter 5 DIsc~~~siorr 68
1 I
Stage 1 / Stage 2 Stage 3
O 5 I O 15 20 25 30 35 40 45
Time (hrs)
Figure 5.4:
These
Weight gain for the synthetic 5000 series alloy at 700°C showing three different oxidation kinetics.
three sections represent difTerent oxidation behaviors occurring at the
sampie surface. Stage one fits a parabolic relationship and is representative of a diffusion
controlled oxidation process. Initial oxidation involves the oxidation of magnesiurn to
fonn MgO. Oxidation continues through the thickening of the initial oxide layer, most
likely through a difision controlled process. As the thickness increases, the oxidation
slows down due to the increased diffusion distance. At some point, the oxide layer cracks.
This is due either to the stresses that developed according to the volume decrease
associated with the Mg to Mg0 reactioo, or to the nucleation of large spinel crystals at
the oxide/metal interface. Once the oxide cracks, Fresh metal is exposed and oxidation
continues in a linear manner. There are two linear regions to the graph. The transition
from one to the other is quite graduai and is most likely due to a change in the surface
area open to oxidation.
X-ray difbction results indicated only the presence of Mg0 as the top oxide.
Generally, spinel has oniy been found to form when an area of magnesium depletion
exists within the sample. Formation of a spinel layer would occur at the oxidelmetal
Cha~ter 5 Discussion 69
interface. Therefore, a layer of Mg0 greater than seceral microns in thickness would
expiain the non-detection through XRD analysis. However, it is possible that not enough
magnesium was consumed through oxidation to cause any areas of depletion and the
subsequent formation of spinel.
Oxidation at 750 and 800 "C
Oxidation rates at 750 and 800°C follow a different initial oxidation behavior.
Initial sections of the weight gain curves do not exhibit a multi-stage oxidation behavior.
Curves are charactenstically linear immediately after the oxidation experiments have
been started (Figure 5.5).
O 2 4 6 8 10
Time (hrs)
Figure 5.5: An example of the initial oxidation period of a 750 and 800°C sample.
This is likely due to the effects that higher temperatures have on the formation of
the initial oxide. This forms so quickly that &acture occurs immediately, and linear
oxidation ensues. It is also possible that the initiation of oxidation occurred before weight
gain measurements actually started. It is believed that the oxide forming at this time is
Mg0 and is limited to the surface of the sample. It is important to note that, although
Chanter 5 Discussion 70
weight gains associated with initial oxidation petiods are low cornpared to the other
stages of oxidation, the oxidation behavior is not considered protective. Given enough
time, continuation of oxidation at the same rate as that observed in the initial periods
would result in the eventual consumption of al1 the magnesium.
5.4.2 Rapid Oxidation
This stage of oxidation was iimited to sampirs o x i d i ~ d ÜL 750 aiid SûO°C. Ai
some point dunng oxidation, the growth rate increased dnrnatically. The inctease in
growth rate is directly related to the amount of oxidation occumng. For al1 samples
tested, the rapid oxidation stage continued until al! of the magnesium was depleted from
the sample. Figure 5.6 shows elemental scans indicating mapesium depletion from the
bulk of the sample.
Figure 5.6: Elemental scans of a) Al and b) Mg of a selected part of a sarnpie oxidized at 800°C.
Some variation existed in the shape of the weight gain curve of the rapid
oxidation stage. Samples tested at 800°C had a more consistent and higher average
weight gain rate than the samples tested at 750°C. It took, on average, approximately 7.5
hours from the initiation of npid oxidation for the 8ûû°C samples to exhaust the
magnesium supply, while the 750°C samples took apprcximately 27 hours. This c m be
explained by undestanding the driving force behind this type of oxidation.
Discussion
Initiation
The fact that oxide clusters are found beneath the oxidizuig surface, as shown in
Figure 4.27% leads to the conclusion that localized areas of elemental magnesium form
beneath the initial oxide layer. Since magnesium is soluble in aluminum in its liquid state,
these localized areas are likely to be gaseous. This statement is strengthened by the fact
that magnesium will vaporize at the temperatures used throughout this study.
Previous studies involving the midation of .Ai!-Mg dloys are uxlear on the
initiation of breakaway oxidation. Some view the phenomenon as being due to nucleation
of large spinel crystals under the primary Mg0 oxide. These large crystals crack the
oxide and allow for continued oxidation through wicking of molten metai to the oxidation
front. A second view is simila. to the one adopted in this study where pockets of
magnesium vapour form below the surface and fracture the initial oxide. Rapid oxidation
ensues through the wicking of molten metal to the oxidizing front. It is possible that the
first viewpoint is true in this study except that it does not explain the presence of so rnany
oxide clusters undemeath the sample surface. In a sirnilar manner, the second hypothesis
is sornewhat unsubstantiated in that if a pocket of magnesium vapour formed, it would be
located dong the top surface of the sample due to the buoyancy forces in effect between a
gas phase and a liquid phase. Intempt experiments that were performed indicated that
initiation took place midway dom, dong the sides of the sample where no pockets of
magnesium vapour should exist.
In reality, it is likely that a combination of the two mechanisms occur
simultaneously. Pockets of gaseous magnesium do exist. Evidence of this can be seen by
the presence of small crystals, which are likely MgO, on the surface of the interrupted
samples at the point of breakaway initiation (Figure 5.7). Perhaps breaking of the oxide is
caused by the reaction of the M g 0 oxide layer with the walls of the alumina crucible to
form spinel crystals that lack the cohesion of the MgO. As the oxide cracks, molten metai
is released through the crack, the pockets of magnesium gas then flow out with the metal
and oxidize in a rapid manner. This would account for both the presence of magnesium
vapour and the initiation of rapid oxidation occurrhg dong the crucible wails. This
theor- also better explains the odd shapes forrned throughout the oxidation experiments.
It is believed that, initially, magnesiurn oxidizes to form MgO. There is a great deal of
Chnpter 5 Discussion 72
turbulence that occuis durinp this oxidation. Since the sample is molten, any
displacement of the oxide layer will affect the underlying metal, accounting for the odd
shapes that are forrned throughout the 750 and 800°C tests. A schematic of the proposed
mechanism is shown in Figure 5.8.
Figure 5.7: The presence of smali crystals (potentially MgO) on the sarnple surhce near the point OF breakaway initiation
Propagation
Continuation of the rapid oxidation stage is different for the 750 and 8ûû°C
sampies. At 8W°C, the curve is not as wavy, indicating a steadier oxidation rate. The
rapid oxidation stage will continue until al1 of the rnagnesium has been oxidized. It can be
seen from the specific weight gain curves, that al1 curves at 750 and 800°C become
horizontal near the 100% Mg exhaustion lirnit. The variance between the maximum
oeneous specitic weight gain values of the curves can be accounted for by the hetero,
distribution of magnesium throughout the samples. It is possible that the magnesium
content from sample to sample varied slightly. A difference of only 0.2% Mg, translates
to a change of approximately 4 1 in potential specific weight gain. Additionally, it is
possible that some alurninum oxidizes to f o m A1203. This would account for the specific
weight gain values that exceed the 100 1 range.
Cha~ter 5 Discussion 73
Sam pie Al203
C ruci ble Surface
Figure 5.8: Schematic of the proposed mechanism responsible for the initiation of brerliaway ouidation. a) Formation of magnesiurn vapour pockets below the top oxide layer, b) Cracking of the oxide layer along the sides of the sarnple due to spinel crystal formation. c) Flow of metal out of crack and release and immediate oridation of magnesium vapour pockets.
Samples oxidized for the fui1 15 hours displayed oxide clusters that were
comprised of two different phases: pure molten alurninum and spinel. The center of the
clusten contained large amounts of porosity. There was more oxide present around the
pore perimetrr indicating that the clusten grew into the metai from the pores.
Continuation of the oxidation of the clusters is best understood by observing samples that
had not consumed al1 the magnesiurn.
Chapter 5 Discussion 74
The clusters that were still in the process of oxidizing were comprised of two
main sections, an inner dense, highiy oxidized region surrounded by a less oxidized area.
The h e r area was always centered around a large amount of porosity. It was found that
the inner area was comprised of spinel and pure duminum while the outer area was
comprised of Mg0 and aluminum. The growth of these clusters continues outwards fiom
the pore. The initial reaction involves magnesiurn and oxygen to form MgO. As the
growth front continues. the metal entrapped in the oxide mabix is devoid of magnesium.
The magnesiurn depletion allows for a secondary reaction to occur. The aluminum reacts
with the Mg0 to form spinel. Since only metal that is far Frorn the oxidizing front is
depleted in magnesium, spinel is located only near the center of the clusters in samples
that have not been hlly oxidized. The outer regions of the cluster continue to react to
form MgO. Clusters observed in samples that had fûlly oxidized contained only spinel.
This is supported by the x-ray difiaction pattern obtained fiom the milled powder s h o w
in Figure 4.28a. It is believed that enough t h e had passed to transform al1 the Mg0 into
spinel.
The oxidation behavior of the synthetic 5000 senes alloy changes greatiy between
700°C and 800°C. From the weight gain curves it is evident that commercial practices
should keep holding and casting temperatures as low as possible in order to minimize the
loss of magnesium through oxidation. Not only will the melt chemistry be affected
through the preferentiai oxidation of magnesiurn, but the amount of aluminurn lost to
entrainment within the dross will aiso increase drastically with increasing temperatures.
5.5 OXIDATION OF COMMERCIAL 356 ALLOY
The commercial 356 dloy contained 0.35 % Mg and therefore was expected to
have substantial weight gains due to the preferred oxidation of magnesium. In the same
way as that used for the 5000 series alloy, the specific weight gain of the 356 alloy has
been plotted in Figure 5.9. Udike the 5000 series aiioy, none of the samples oxidized
approach the magnesium exhaustion limit. This might be due to some small loss of
magnesium by evaporation. Additionally, the c w e s do not have the same shape as the
Cha~ter 5 Discussion 75
5000 series alloy. Oxidation is characterized by a steady growth rate that continually
decreases as time progresses. No breakaway oxidation stage exists throughout the full 45
hours for any of the sarnples tested.
Time (hrs)
Figure 5.9: Specific weight gains of the commercial 356 alloy.
The driving force for the rapid oxidation stage in the synthetic 5000 senes alloy is
believed to be the buildup of magnesiun vapor beneath the initial oxide layer and
subsequent release of this vapour due to cracking of the initial oxide through spinel
formation. However, the commercial 356 ailoy has less than one-tenth the magnesium
content of the synthetic 5000 series ailoy and therefore cannot build up pockets of
rnagnesiurn. Oxidztion therefore continues in a similar manner to the induction period of
the 5000 series alloy. Growth is controlled by the reaction of magnesium to form MgO.
In a region of depleted magnesium, Mg0 will react with aluminum to form spinel. The
decrease in weight gain rate with time is representative of the decreasing magnesium
content. Less Mg0 is forming at longer times and therefore weight gains are due more to
the reaction of Mg0 with aluminurn to form spinel than with the formation of MgO. The
speed of these reactions changes gradually with the decreasing concentration of
magnesium in the melt. As more magnesium is consumed, less is available for reaction.
Discussion
In addition, as time progresses, less Mg0 is available to react to form spinel. The weight
gain rate therefore decreases with time. Had the experiments continued for another 45
hours, it is likely that the oxidation rate would have slowed down to a negligible value
due to the complete exhaustion of magnesium and MgO.
The effect of temperature on the oxidation of the 356 alloy is evident. Weight
gains for the 800°C samples are approximately fours times greater than the 700°C
samples. In a similar manner to the 5000 series alloy, the difference between the 700°C
samples and 750°C samples is much greater than the difference between the 750°C and
800°C samples. It is believed that the temperature not only increases the kinetics of the
reactions involved during oxidation, but also increases the availability of magnesiurn to
zones open to oxidation. The viscosity of molten aluminum alloys decreases with
increasing temperature, and it is reasonable to assume that more molten metal can
penetrate through to the oxidizing fiont. The efTect of temperature is therefore twofold.
Higher temperatures increase reaction kinetics and the supply of fiesh, unoxidized metal
through cracks in the oxide.
5.6 EFFECTS OF STRONTIUM ON COMMERCIAL 356 ALLOY
Strontium was added to the 356 alloy in order to observe the effects that it would
have on the oxidation behavior of the alloy. The difference in weight gains with the
addition of strontium was drarnatic. Weight gains decreased dramatically, by
approximately ten times. The duIl gray appearance of a heavily oxidized sample was not
observed on the surface of any of the strontium containing samples.
Upon examination of the samples through microscopy, smail, localized areas of
oxidation were observed on al1 samples. The addition of strontium effectively reduced the
oxidation by establishing a more stable or stronger top oxide. The amount of strontium
added to melt was relatively small (250 ppm) and therefore is not abundant enough to
form a thick complete oxide layer on its own.
Chanter 5 Discussion 77
Strontium c a . oxidize to form SrO. If this oxide had formed over the entire
surface and al1 250 ppm had been consurned in the formation of this oxide layer, then the
resulting thickness of the oxide would be approximately 3 p. However, the complete
consumption of strontium did not occur and no oxide was observed through cross-
sectional analysis. This supports the presence of an oxide layer thinner than 1 p. Post
test examination of the microstructure of the strontium containhg 356 alloy revealed the
yresence of scne rern-ining strontitm.. .4!thcug!h qmtirr?ir.e mdysis of the renahiiig
strontium could not be perforrned due to the porous nature of the samples d e r testing, it
is evident that not al1 the strontium was consurned during the oxidation test.
In order to examine the effect of strontium on the oxidation of the commercial
356 alloy, it is necessary to compare oxidation curves of alloys with and without
strontium. By tnincating the weight gain curves for the two different alloys at 2.5 hours it
is possible to see that strontium increases the initial oxidation rate (Figure 5.10). C w e s
for the strontium containing alloy were fit to a logarithmic relationship.
O 0.5 1 1.5 2 2.5 Time (hm)
Figure 5.10: Cornparison of the initial oxidation periods of the commercial 356 ailoys.
The difference between the curves of the MO 356 alloys is evidence that strontium
plays a role in the oxidation of Al-Mg alloys. There are two possible ways in which the
strontium can affect the oxidation. Fintly, it can preferentially oxidize to form an oxide
of its own, such as Sr0 or SrO2. Secondly, it can combine with the other elements present
Chapier 5 Discussion 78
to form a more complex oxide such as SrA1204. In either case, the oxide that forms is
likely to be very thin due to the srnail amount of strontium present in the alIoy.
X-ray difiction of the strontium containing 356 alloy revealed intense peaks
representing aluminurn and silicon. However, some small, low angle peaks were
observed in the pattern. Low angle peaks indicate a large unit-ce11 dimension. Although
no conclusive match was made for the additional smail peaks, the best candidates were
AlzSilSr and SrA407. Both of the compounds had partial matches where some o f the
peaks were present and others were rnissing (Figure 5.1 1). Since the signal of the
materials is so weak, the smaller peaks in their respective spectra are possibly lost in the
background. In addition, any texturing of the oxide, or preferred orientation, might resuit
in the absence of some peaks.
I O 15 20 25 30 35 40 45 50 55 60 Angle ('20)
Figure 5.11:X-ray diffraction pattern showing srnail peaks matching strontium- containing phases.
Microchemical analysis of approximately the top 3 pn of the 356 alloy containing
strontium contradicted the x-ray difiaction results. The presence of magnesium was
detected dong with peaks related to silicon, strontium, oxygen, and aluminum. Peaks for
copper, gold, and palladium are due to the coating placed on the sample and should be
Chapter 5 Discussion 79
ignored. What is evidrnt from the spectrurn shown in Appendix 1 is that the aluminum
peak is very weak and the Mg peak is very intense. Perhaps, examination through EDS
detects two different oxide layers. The top layer, detected through x-ny diffraction is a
very thin strontium containing oxide, while undemeath this first oxide is a thicker layer
of MgO.
An outline of this hypothesis is as follows: the initial oxide that forms is an AI-Si-
Sr-O specirs. Tliis a i d e k m i s ~ : i i î ~ m ~ l : f quiskly asras thc iûp of thc inoitcn sürhcc.
The oxide is extrernely protective, thereby limiting further oxidation. Further formation
of this top oxide is lirnited due to the small arnount or strontium. Additional oxidation
occurs through the slow diffusion of ions through this initial layer and reaction to form
MgO. The diftïculty with this hypothesis is that if two separate oxides fonnrd, the weight
gain curves would likely be discontinuous when the transition from one oxide to the other
occurs. This is not observed on the weight gain curves seen in Figure 4.15.
A second hypothesis proposes that M g 0 foms on the surface of the sarnple. The
M g 0 layer is stabilized by a solid solution of strontium within the crystal structure. Both
Mg0 and Sr0 have the same electron valence number and f o m a similar rocksalt ceIl
structure, and therefore a solid solution of strontium within Mg0 is quite plausible.
However, this hynothesis is unlikely due to the fact that strontium should not stabilize
MgO. The P-B ratio of strontium (0.69) is less than that of magnesium (0.83). Therefore,
as more Sr atoms replace the Mg atoms in the crystal structure, the P-B ratio converts to a
value further form unity, and hence, to an oxide with less protective properties. In
addition, this hypothesis tàils to explain the presence of small, low angle peaks in the x-
n y diffraction pattern.
A third possibility exists. Similar to the case of the synthetic 413 alloy, an initial
oxide exists, except in this case it is MgO. This oxide is fomed at room tempenture
before the heat ing cycle has begun. When the erperiment is started, strontium oxidizes
preferentially with aluminum or silicon to form SrA140: or SrSiO,. The formation of this
oxide occurs below the initial Mg0 layer formed at room tempenture. It is protective and
prevents further oxidation, except for the occasional localized eruption. These eniptions
are caused by small, localized build-ups of magnesium vapour below the surface oxide.
Since a protective oxidr forrns, a thickening of the initial M g 0 layer is prevented, dong
Chavter 5 Discussion 80
with consumption of magnesium. Since no magnesium is being depleted, gas builds up at
the rnetalloxide interface. This does not occur in the commercial 356 ailoy because the
continual thickening of the top Mg0 oxide layer prevents the build up magnesiurn vapour
at the surface of the sample. This hypothesis explains the fact that the weight gain curves
follow a different kinetic law through the formation of a diEerent oxide. Secondly, it
explains the presence of Mg0 indicated in the EDS anaiysis by assurning an initial Mg0
layer. Thirdly, it explains the presence of small. low angle peaks in the XRD analysis by
assurning the formation of a strontium containing oxide with a large unit-ce11 dimension.
However, it is important to question whether the oxide formed on the strontium
containing 356 alloy is the same as that formed on the strontium containing synthetic 4 13
alloy. If that is the case, their weight gain curves should exhibit the same oxidation
behavior. Perhaps, since the initial oxide on the 413 alloy is A1203, a concentration
gradient is formed as outlined by the Wagner mechanism shown in Figure 3.3. This
would allow for continual growth through the formation of numerous oxides with
differing concentrations of alurninum, silicon, strontium, and oxygen.
6. CONCLUSIONS
Oxidation experirnents were conducted on six different alloys: pure alurninum,
synthetic 41 3 alloy, synthetic 4 13 alloy containing 350 pprn of strontium, synthetic 5000
senes alloy, commercial 356 ailoy, and commercial 356 alloy containing 250 ppm
strontium at temperatures ranging fiom 700 to 800°C.
It has been shown thz? the addition of 12% silicon to aluminum will increase the
oxidation behavior through the formation of a-alumina oxide nodules. Silicon had no
effect on the oxidation behavior at lower temperatures.
The addition of strontium to the synthetic 413 alloy effectively prevented the
formation of a-dumina oxide nodules fiom forming at 800°C. This occurs through
the formation of a strontium containing species at the surface of the sample. It was
further concluded that strontium did not accelerate the initial oxidation behavior of
the synthetic 4 13 alloy.
Magnesium will preferentiaily oxidize over aiuminum to form firstly, MgO, and
secondly spinel. Spinel formation occurs due to a depletion of magnesium widiin the
surroundhg area Oxidation of Al-3% Mg alloys will only stop oxidizing once al1 the
magnesium has been exhausted.
Temperature had a large effect on the oxidation behavior of the synthetic 5000 senes
alioy. Higher temperatures not only increase the reaction kinetics but dso reduce the
induction period and time required to oxidize the magnesium within the alloy.
Chapfer 6 Conclusions 82
Breakaway oxidation in the synthetic 5000 senes alloy was initiated through a
combination of two effects. a) The formation of spinel crystals at the surface oxide
and subsequent cracking of the oxide. b) The development of gaseous magnesium
pockets below the surface which once releaçed through cracking of the top oxide
cause sharp increases in weight gains. This oxidation behavior causes the formation
of large clusters of oxides within the bulk of the sarnple.
Cornmercil! 726 d ! ~ y r do not exhibit breakaway oxidation due to its lower
magnesium content. Build-up of magnesium vapor is prevented through the continual
growth of the top oxide layer.
The addition of strontium to commercial 356 alloys results in a drastic decrease in
oxidation of the ailoy. It is proposed that a strontium-containing phase forms
undemeath the initial Mg0 layer that was formed at room temperature. This oxide is
protective and very thin, causing weight gains to remain low.
7. SUGGESTED FUTURE WORK It is evident that the details of how strontium interacts with a l h u m alloys are
stiil uncertain. Additional work, focusing on the identification of the strontium containing
oxides would shed some light on exactly what is occurring during oxidation. In order to
accomplish this goal, larger scale experiments should be conducted where enough oxide
is f o n e d to dlow for proper XRD analysis. Care should be taken as to the type of
crucible used for these experiments. A high purity carbon crucible might prevent
oxidation dong the sides of test sarnples.
Additionally, experiments on Al-Mg alloys containing strontium should dso be
perfomed to determine if the presence of silicon is necessary in order to prevent
magnesium oxidation.
Finally, a study of the effect of humidity on the oxidation behaviour of aluminurn
alloys containing strontium should be conducted. The change in hurnidity associated with
the change in seasons and location is thought to have an effect on the oxidation behaviour
of aluminum alloys. Simulation of various humidities can be accomplished through the
introduction of predetermined air humidity contents during oxidation experiments.
APPENDIX 1
EDS Analysis Results
a) Synthetic 413 alloy containing 350 ppm of strontium: Area analysis of the surface.
b) Commercial 356 aiioy containing 250 ppm of strontium: Area analysis of the surface.
c) Synthetic 5000 series alioy: Area analysis of base metal region.
d) Synthetic 5000 series alloy: Spot analysis of oxide in the inner section of a cluster.
e) Synthetic 5000 series alloy: Spot analysis of metal channels in the inner section of a cluster.
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