The Effect of Processing Parameters and Alloy
Composition on the Microstructure Formation and
Quality of DC Cast Aluminium Alloys
Majed M. R. Jaradeh
Doctoral Thesis
In
Materials Processing
Department of Material Science and Engineering
School of Industrial Engineering and Management
KTH-Royal Institute of Technology
SE-10044 Stockholm, Sweden
Department of Engineering, Physics and Mathematics
Mid Sweden University
SE-851 70 Sundsvall, Sweden
Akademisk avhandling som med tillstånd av Kungliga Tekniska Högskolan (KTH)
framlägges till offentlig granskning för avläggande av Teknologie Doktorsexamen,
Todsdagen den 14 Dec 2006, kl 10.15, sal F3, Lindstedtsvägen 26, KTH, Stockholm.
i
The Effect of Processing Parameters and Alloy Composition on the
Microstructure Formation and Quality of DC Cast Aluminium Alloys
Majed Jaradeh
Department of Material Science, Royal Institute of Technology
SE-100 44 Stockholm, Sweden
Abstract
The objective of this research is to increase the understanding of the solidification behaviour
of some industrially important wrought aluminium alloys. The investigation methods range
from direct investigations of as-cast ingots to laboratory-scale techniques in which ingot
casting is simulated. The methods span from directional solidification at different cooling
rates to more fundamental and controlled techniques such as DTA and DSC. The
microstructure characteristics of the castings have been investigated by optical and Scanning
Electron microscopy. Hardness tests were used to evaluate the mechanical properties.
The effects of adding alloying elements to 3XXX and 6XXX aluminium alloys have been
studied with special focus on the effects of Zn, Cu, Si and Ti. These elements influence the
strength and corrosion properties, which are important for the performance of final
components of these alloys.
Solidification studies of 0-5wt% Zn additions to 3003 alloys showed that the most important
effect on the microstructure was noticed at 2.5 wt% Zn, where the structure was fine, and the
hardness had a maximum. Si addition to a level of about 2% gave a finer structure, having a
relatively large fraction of eutectic structure, however, it also gave a long solidification
interval. The addition of small amounts of Cu, 0.35 and 1.0 wt%, showed a beneficial effect
on the hardness.
Differences have been observed in the ingot surface microstructures of 6xxx billets with
different Mg and Si ratios. Excess Si compositions showed a coarser grain structure and more
precipitations with possible negative implications for surface defect formation during DC
casting.
The comparison of alloys of different Ti content showed that the addition of titanium to a
level of about 0.15 wt% gave a coarser grain structure than alloys with a normal Ti content for
grain refinement, i.e. < 0.02 wt%, although a better corrosion resistance can be obtained at
higher Ti contents. The larger grain size results in crack sensitivity during DC casting.
A macroscopic etching technique was developed, based on a NaOH solution, and used in
inclusion assessment along DC cast billets. Good quantitative data with respect to the size and
spatial distribution of inclusions were obtained. The results from studied billets reveal a
decreasing number of inclusions going from bottom to top, and the presence of a ring-shaped
distribution of a large number of small defects in the beginning of the casting.
The present study shows how composition modifications, i.e. additions of certain amounts of
alloying elements to the 3xxx and 6xxx Al alloys, significantly change the microstructures of
the materials, its castability, and consequently its mechanical properties.
Keywords: Aluminium wrought alloys, AA3xxx and 6xxx, Direct Chill Casting,
Unidirectional Solidification, Bridgman technique, Process parameters, Microstructure,
Composition modification, Thermal analysis, Inclusions, Metallographical investigations.
ii
Acknowledgments
This work has been carried out at Mid Sweden University, at the Department of
Engineering, Physics and Mathematics. I would like to express my sincere
appreciation to my professors, Torbjörn Carlberg, the main thesis supervisor and
Hasse Fredriksson (co-supervisor), for their scientific guidance, encouragement and
kindness throughout my graduate studies and the completion of this dissertation. The
great inspiration and confidence that they have given me make me believe that I can
continue my academic dream. I am also grateful to the staff members of the Sundsvall
branch of the Department of Engineering, Physics and Mathematics, at Mid Sweden
University and the Department of Materials processing at the Royal Institute of
Technology in Stockholm, especially Advenit Makaya. The financial support from
EU, Objective 1, has been of great importance for my investigation, and my coworkers and I would also like to thank Kubikenborg Aluminium AB for their support
and provision of the sample materials.
Last, I would like to thank my family and friends who have motivated and supported
me through the completion of this dissertation.
Sweden, December 2006
Majed M. R. Jaradeh
iii
List of papers
The thesis includes the following supplements:
Supplement 1
Analysis of solidification in a Bridgman furnace as a simulation of DC casting of
aluminium alloy slabs.
Majed Jaradeh and Torbjörn Carlberg.
Accepted for publication in Materials Science and Technology, 2006
Supplement 2
Solidification studies of 6xxx alloys with different Mg and Si contents.
Hu Jin, Majed Jaradeh and Torbjörn Carlberg.
Proceedings, Light Metals. TMS, San Francisco, California, USA, Feb. 2005, pp. 1039-1044.
Supplement 3
Directional solidification of type AA3003 alloys with addition of Zn and Si.
Majed Jaradeh and Torbjörn Carlberg.
Proceedings, Solidification of Aluminum Alloys. TMS, Charlotte, North Carolina, USA.
March 2004, pp. 53-62.
Supplement 4
DTA and DSC Studies of aluminium 3003 alloys with Zn and Cu additions.
Majed Jaradeh and Torbjörn Carlberg.
Submitted to Metallurgical and Materials Transactions A. 2006
Supplement 5
Effect of titanium additions on the microstructure of DC-cast aluminium alloys.
Majed Jaradeh and Torbjörn Carlberg.
Materials Science and Engineering A. Vol. 413-414, 15 December 2005, pp. 277-282.
Supplement 6
Method Developed for Quantitative Assessment of Inclusions in Aluminium Billets.
Majed Jaradeh and Torbjörn Carlberg.
To be published in Proceedings, Light Metals, TMS. Orlando, Florida, USA. Feb. 2007.
Supplement 7
Solidification Studies of Automotive Heat Exchanger Materials.
Torbjörn Carlberg, Majed Jaradeh and H. Kamgou Kamga.
JOM. Vol. 58, no. 11. Nov, 2006.pp 56-61
iv
TABLE OF CONTENTS
Abstract.......................................................................................................................... ii
Acknowledgments ........................................................................................................iii
List of papers ................................................................................................................ iv
TABLE OF CONTENTS ............................................................................................. v
Chapter 1 ....................................................................................................................... 6
1
Introduction and Background of Thesis ............................................................ 6
1.1 Direct Chill Casting of Aluminium Alloys...................................................... 6
1.2 Aluminium Alloys and Effects of Alloying Elements..................................... 9
1.3 Aluminium Melts Cleanliness and Treatment ............................................... 11
1.4 The Scope of the Thesis ................................................................................. 14
1.5 References ...................................................................................................... 15
Chapter 2 ..................................................................................................................... 19
2
Experimental methods ....................................................................................... 19
2.1 Materials......................................................................................................... 19
2.2 Experimental Techniques............................................................................... 19
Chapter 3 ..................................................................................................................... 28
3
Results and discussion........................................................................................ 28
3.1 Results of thermal conditions in the Bridgman furnace................................. 28
3.2 Results of different Mg and Si contents on solidification of 6XXX alloys... 32
3.3 Results of additions of Zn and Si on the solidification of 3003 alloys.......... 35
3.4 Results of DTA and DSC studies of 3003 alloys containing Zn and Cu....... 37
3.5 Results of additions of excess titanium.......................................................... 41
3.6 Results of assessment of inclusions in aluminium billets.............................. 43
4
General Conclusions of Thesis .......................................................................... 46
v
Chapter 1
1
Introduction and Background of Thesis
Aluminium and its alloys have excellent properties: they are light in weight, corrosionresistant, good electrical and thermal conductors and relatively easy to produce. They
find their main uses in the transportation, construction, packaging and electrical
industry. Aluminium can be produced via two different routes: primary aluminium
production from ore, and recycling from process scrap and used aluminium products.
The production of primary aluminium consists of three steps: bauxite mining, alumina
production and electrolysis. The bauxite is refined into alumina (Al2O3) by the Bayer
process and in reduction plants primary aluminium is extracted from alumina by the
Hall-Heroult electrolytic process [1-3].
For all aluminium alloy components, the casting process plays an important role in
controlling the properties of the final products. Semicontinuous direct-chill (DC)
casting is currently the most widely used process in the production of aluminium
alloys [4,5]. This casting technique is commonly used to produce ingots that undergo
subsequent processing. Ingots of rectangular cross-section are commonly rolled to
sheet, foil and plate; cylindrical ingots are extruded to form rods, bars and wires [6].
The casting method was developed independently by VAW (Germany) and Alcoa
(USA), based on the patent of Roth and W. T. Ennor in the early 1930s and 1940s
respectively [7,8]. A brief detail of the process description and DC machine casting is
reported below.
1.1 Direct Chill Casting of Aluminium Alloys
In the semicontinuous direct chill (DC) casting process, molten aluminium is poured
from the furnace through a troughing system into a mould table containing moulds
with a cross-sectional shape of the desired ingot, which may be rectangular for rolling,
round for extrusion billets. A starting head, mounted to a hydraulic ram, forms a sort
of false bottom to the mould. When the metal in the mould reaches a definite height,
the bottom block is lowered at a constant rate equivalent to the liquid metal flow. A
skin solidifies on the first contact of the melt with the mould wall, forming a shell to
hold the liquid metal. The direct chill is applied just below the mould exit by spraying
cooling water onto the ingot to cool the surface and further solidify the metal. The
water-cooled mould and the water sprayed on the hot ingot cause a very rapid
solidification. Casting speeds normally vary with the ingot size and alloy type and are
typically in the 1-3 mm/s range (1. 10-3 to 3. 10-3 ms-1) [4]. The cooling rate varies in
relation to the mould walls and is approximately 0.5 oC/s in the ingot centre to
approximately 20 oC/s near the ingot surface. The thermal gradient is 0.3-5.0 K/mm,
and the pouring temperature of the aluminium melt alloys is near 690-725 oC. The
water flow rate is ramped up in the beginning but constant during steady state castings.
The DC casting apparatus of rectangular cross-sections used for rolling ingots, shown
in Fig.1.1, is composed of a water-cooled rectangular mould, made of a copper or
aluminium alloy, with a plug or ingot stool resting on a casting table.
Figure 1.1: Schematic of direct-chill (DC) casting machine.
Recent developments in the DC casting include the electromagnetic casting, the hottop system and the air slip process. The air slip mould was developed specifically for
production of billets. The process involves the formation of a gas barrier between the
molten metal and the graphite-lined mould [5,9]. The benefits gained by industries
from using the mentioned new variants of the process over the basic DC cast process
are the improvement of the surface microstructure and minimization of surface defects.
Although the DC casting technology has been used for a long time, several problems
are still unsolved, particularly in relation to the demand for increased productivity and
the reduction of defects (e.g. ingot cracking, shrinkage, porosity, inclusions, etc.).
From a productivity standpoint, the casthouse would preferably produce the largest
possible ingots at the highest possible casting speeds, but reality intervenes, with ingot
size and aspect ratio (ingot thickness/ingot width) as well as casting speed restricted by
the occurrence of casting defects, such as cracks and hot tears [10-14].
The quality and the commercial applications of the aluminium cast products depend, to
an important extent, on their microstructure. The microstructure variables of interest in
aluminium alloy castings are Dendrite Arm Spacing (DAS), grain size and type,
amount and morphology of intermetallic particles. Each of these variables determines
the mechanical properties of the part. The values of the variables in the as-cast part
depend on the solidification characteristics during casting, such as cooling and
solidification rates, the alloy composition, additives such as modifiers and grain
refiners, etc. [15,16].
A characteristic feature of the process of ingot casting of aluminium and its alloys by a
semicontinuous method is that, the solidification parameters, which influence the
solidified microstructure are not constant across the thickness of the ingot, i.e. each
elementary volume of the cast metal is solidified with different rate. Its magnitude
depends on the parameters used in the casting process, (e.g. the temperature of the
melt, the casting speed, the amount of cooling water), on the position in relation to the
mould walls and the axis of the ingot being cast. It means that this phenomenon is
critical in the case of the casting of large ingots of rectangular cross-section, which is
used as the charge to make rolled products. Wan et al., [17] measured the solidification
rates of DC cast ingots by in situ thermocouple tests. The solidification rate during
cooling was observed to be faster on the surface of the ingot and was observed to
decrease towards the centre of the ingot. This change in the solidification rate during
cooling causes the change in microstructure, which results in a variation of mechanical
properties at different locations of the ingot.
In direct chill (DC) casting of aluminium and its alloys, it is a well-established practice
to add a grain refiner to promote the formation of a fine, equiaxed grain structure, to
improve many properties of the as-cast metal and to facilitate subsequent processing
[18-20]. Grain refinement allows a more uniform distribution of alloying elements,
and it reduces porosity and the risk of defects especially cracks in the cast structures.
Grain refinement by inoculation is commonly achieved by the addition of a master
alloy to the melt, typically in the form of a rod or a waffle. Various types of master
alloys can be used. The most common master alloys are based on the Al-Ti-B system,
(e.g. Al-5% Ti-1% B). The amount that needs to be added is calculated on the basis of
a concentration of 0.02 wt.% Ti (with respect to the weight of the molten metal to be
treated). Typical addition levels of master alloys are between 0.5 and 2 Kg ton-1 of
melt in DC ingots [21]. Much work has been directed to understand the mechanisms of
grain refinements [21-25]. Of all the mechanisms proposed so far, duplex nucleation
theory is the most recent, first proposed by Mohanty and Gruzleski [26] and further
developed by Schumacher and Greer [27-29]. This theory explains the mechanism of
grain refinement of Al alloys by the addition of AlTiB master alloys, which suggests
that TiB2 particles need to be coated with a thin layer of Al3Ti to be effective nucleant
particles. A second important effect is the grain growth restriction, i.e. when alloying
elements restrict the growth of aluminium dendrites and form an undercooled melt
zone [30]. Ti is the most potent growth restriction element among all other alloying
element in aluminium alloys. Both the partitioning of solute elements and the added
nucleant particles are now thought to affect the grain refinement process.
Minimizing casting defects and improving the productivity of the DC-casting
processes through a better understanding of the mechanisms involved during the
casting and the solidification are the focal points of efforts worldwide. However, the
investigation of an industrial scale ingot is expensive, time-consuming and tedious, so
it is becoming increasingly important to study solidification process and
microstructure formation within the laboratory. A number of small-scale casting
techniques have been developed to minimize the need for the casting of full size ingots
during research on DC-casting phenomena and the development of new alloys. These
experimental techniques [21,31] either mimic solidification conditions within the ingot
(e.g. Direct-chill, DC simulator) [32,33] or allow solidification to occur under
controlled conditions, so that the fundamentals of solidification can be investigated
(e.g. Bridgman technique). These two directional solidification techniques were used
in the course of this work to develop our understanding of the microstructure
formation, and they are fully described in chapter 2, section 2.2.1. Both techniques
have also been used previously to investigate the formation of intermetallic-phase
selection [34-37], and fundamental theory associated with grain refinements in
wrought aluminium alloys [38,39].
Other types of experiments available to trace changes in the solidification and
precipitation of aluminium alloys involve the use of thermal analysis techniques. Two
types of thermal analysis methods have been used in this research: differential thermal
analysis (DTA) and differential scanning calorimetry (DSC). These devices are well
documented in literature and widely used to calculate material properties such as latent
heat and specific heat and to determine the reaction kinetics during phase
transformations [40-42]. Although the development of DTA and DSC instruments
made it possible to determine values faster, more conveniently and at a lower cost,
these techniques can be less accurate to investigate the solidification of DC-casting
conditions and to predict commercial DC-cast microstructure. The solidification of a
DC-cast ingot is complex, and normally it does not occur under equilibrium
conditions. The techniques mentioned were designed to determine equilibrium
properties and are limited to small sample volumes. Further details of the experiments
and the analysis methods involving the DTA and DSC techniques applied in this
research are detailed in chapter 2, section 2.2.2.
The aluminium can be tailored to the applications required by modifying the
composition of its alloys, i.e. through the addition of different kinds and quantities of
alloying elements. Many metals can be added to aluminium, according to the product
demand. Manganese (Mn), silicon (Si), magnesium (Mg), copper (Cu), iron (Fe), zinc
(Zn) and titanium (Ti) are among the most commonly used alloying elements. The
different alloying elements have different functions because of their different solid
solubilities in aluminium. Some elements are used as solid-solution strengtheners,
others form desirable intermetallic compounds and therefore induce different effects
[43].
In the next section the purposes of adding different alloying elements are reviewed as
well as their effects on the microstructures and properties important for their
applications.
1.2 Aluminium Alloys and Effects of Alloying Elements
In general, pure aluminium has a very low yield strength of about 7 Mpa, i.e. it is too
soft and ductile to be used in structural and industrial applications, and therefore it
needs to be hardened and strengthened. One common way to strengthen a metal is to
add alloying elements that cause both precipitation and solution hardening, which
results in a high strength-to-weight ratio. The amount of useful alloying elements,
which can be incorporated, is controlled by their solid solubility and the cooling rate
achieved during casting. Since there are no allotropic phase transformations in
aluminium, much of the control of the microstructure and properties relies on
precipitation reactions. The solubility of solute in the matrix is therefore of
importance. Up to 70 wt.% of zinc can dissolve in aluminium, followed by 17.4 wt.%
magnesium, 5.65 wt.% copper, 1.82wt% manganese and 1.65wt% silicon [44].
Certain properties of aluminium are largely independent of the content of alloying
elements and the processing chain. Examples are the modulus of elasticity of common
wrought alloys, which is important for many applications, E = 70 GPa, the density, S =
2700 kg/m3, and the thermal expansion coefficient, 24 x 10-6 /K. Most other properties,
for example strength and corrosion resistance, are very sensitive to the composition
and microstructure of the material [45].
Basically, two types of alloys have been studied for this thesis; the AA3xxx (Mn rich)
and AA6xxx (Al-Mg-Si) series. The aluminium alloy 3003 type is a preferred alloy for
heat exchanger materials in car radiators or air conditioning units [46, 47]. This type of
alloys has the advantage of being non-heat treatable, i.e. they do not require any postbrazing ageing treatment. In this type of alloys, Mn is always present at slightly above
1%. It improves the strength, both by solution and dispersoid hardening, at room
temperature and at brazing temperature, and it also improves the corrosion resistance
and the castability. Common modifications or enhancements in the mechanical
properties of the alloy, i.e. the improvement of strength and enhanced corrosion
resistance are made through the additions of certain alloying elements, such as Zn, Si
and Cu. Additions of different amounts of zinc to the 3003 alloy have been studied, as
this element has a large solubility in aluminium and thus, it can improve the strength,
by solution hardening both at room temperature and at brazing temperature. Another
interesting aspect of Zn alloying is that it opens up the possibility to give corrosion
protection to selected parts of the heat exchangers. Zinc additions lower the corrosion
potential, and can thus be used in part that can act as sacrificial anodes, while other
parts are being protected [48]. Similarly, copper has an electrochemical effect on the
material and alters the corrosion resistance [49]. In addition, Cu contributes to the
solid solution strengthening of the material. Silicon additions are also of interest due to
the low solubility of this element and the potential for the precipitation of virtually
pure silicon and of intermetallic particles of silicon within the matrix, which increases
the hardness and abrasion resistance of the metal. Si also has a low density (2.34 g cm3
), which may be an advantage to keep the overall weight of the component low. In
addition, Si also increases the amount of the AlFeSi phases, which has a positive effect
on mechanical properties. In brazing applications, the most important use of Si is in
clad materials, i.e. as a filling metal during brazing, which has been rolled onto the
core material. Here, the compositions are from 7 to 11% in order to lower the melting
point sufficiently to facilitate the brazing.
The AA6xxx series alloys are very popular in the extrusion industry and are generally
known as high-speed extrusion alloys. This type of alloys is increasingly attractive to
industries such as transportation because of their high specific strength. Alloys in the
6xxx series contain silicon and magnesium approximately in the proportions required
to form magnesium silicide (Mg2Si). The formation of (Mg,Si) particles is facilitated
when the Mg and Si contents are high. The proper ratio for Mg2Si is Mg/Si= 1.73, but
this is impossible to achieve with ordinary operating tolerances; thus, most alloys have
either magnesium or silicon excess. Magnesium excess leads to better corrosion
resistance but lower strength and formability [50]; silicon excess produces higher
strength without loss of formability and weldability [50,51], but there is some
tendency to intergranular corrosion [51]. In the thesis, the influence of alloy
compositions is studied, particularly the balance of Mg and Si concentrations on the
solidification microstructure and ingot surface defects.
Another alloy addition, which is common in many aluminium alloys, is the titanium.
Ti may be added as an alloying element for its grain refinement effect or as a grain
refining agent of Al-Ti-B master alloys. The maximum content of Ti that can usefully
be incorporated in aluminium alloy ingots, cast by normal techniques, is of the order of
<0.02 wt.%, depending on the alloy and grain refinement technique. On the other
hand, if Ti is added in excess of 0.1%, it is believed that improved corrosion resistance
can be obtained [52].
1.3 Aluminium Melts Cleanliness and Treatment
One of the most serious problems in aluminium casting is that cast structures are
usually associated with inclusions. The detrimental effects of the presence of
inclusions are well known. The presence of inclusions can results in increased gas
porosity, poor surface finish, pinholes in thin foils, decreased fluidity, reduced
corrosion resistance, loss of mechanical properties, especially fatigue and ductility,
poor machinability and excessive tool wear, etc [53-57]. Although much work has
been reported on the study of inclusions, the problem of measuring metal cleanliness
through a sensitive quantitative method still exists. The main problem concerning
determining the inclusions size distribution is that they are heterogeneously distributed
in the melt, and that their concentrations are extremely low.
In the following, the main types of inclusions, the removal techniques during the metal
casting of aluminium, and the methods of assessment are reviewed.
Inclusion is defined as any exogenous solid or liquid phase present in molten
aluminium. The types of inclusions found in aluminium are numerous and complex,
and their formation depend primarily on the alloy composition, the melt treatment
practice, and the casting technology. There is three important indices when defining
melt cleanliness: the composition of the inclusions, the particle size of the inclusions,
and the size distribution of the inclusions.
1.3.1 Inclusions: Types and Sources
The inclusions commonly found in cast aluminium alloys include Alumina (Al2O3
particles and films), titanium diborides (TiB2 particles), magnesia (MgO particle), and
magnesium aluminates spinels (MgAl2O4 particles and films) [58]. Inclusions in
aluminium alloys may be from two basic sources [58-60]: the inclusions occurring
from outside the melt, called exogenous, and include refractory particles resulting
from the degeneration of crucibles or furnace linings. They may appear in several
forms with sizes from microns to millimetres. The indigenous inclusions, i.e. those,
which arise from within the melt, may be of several types and are usually only microns
thick (but may extend to several millimetres). The size at which an inclusion becomes
potentially troublesome of course varies with the end-use application. It has been
shown in [61] that the production of quality casting necessitates that the inclusion
contents should be of the order of several volume parts per billion at the most and that
the particle size of the population is less than 50μm. In general, for critical
applications, any inclusion larger than 10-20 micrometers in diameters is considered to
be deleterious, although even smaller inclusions can cause problems if present in
sufficient quantity [54, 57,62].
In the metal casting industry, the treatment and cleanliness of the molten metal is of
great importance. A number of commercially accepted melt treatment techniques are
being used by aluminium foundries to remove and separate inclusions from the molten
aluminium alloy prior to casting. They include various methods of fluxing [63],
degassing, [64,65] and filtration [66-68] in the furnaces and in the gating system.
1.3.2 The Formation of Inclusion
The formation of aluminium oxide inclusions during aluminium casting is inevitable,
since aluminium is very reactive and readily forms a stable oxide on the liquid surface
by reacting with air or moisture. The rate of oxidation increases with temperature and
is substantially greater in molten than in solid aluminium.
While the oxide coatings or oxide skins that initially form on the surface of molten
aluminium are highly protective and self-limiting, any agitation or turbulence in the
treatment and handling of the melt increases the risk of oxide entrainment and
immediate reformation of additional oxides. Oxide concentration can increase when
alloying additions are stirred into the melt, when reactive elements and compounds are
immersed, when the metal is transferred from the conventional metal treatment system
to the casting stations, when the metal is drawn for pouring, and when the metal is
poured into the mould cavity.
The problem of removing aluminium oxides depends on the almost identical density of
the inclusions and the melt. For this reason, the inclusions do not rise to the surface
and tend to remain inside the metal. However, Al2O3 inclusions become buoyant as a
result of absorbed gases, and rise to the melt surface [69].
Magnesium oxide (MgO) and magnesium aluminates spinel (MgAl2O4) inclusions
occur in most aluminium alloy containing magnesium. Since MgO and MgAl2O4 have
lower free formation energy than Al2O3, such inclusions tend to form preferentially,
particularly when the Mg content of the alloy is greater than 0.005 wt.% [70,71].
The addition of AlTiB grain refiner material often results in 1-3 micron titanium
diboride (TiB2) particles. Such small particles individually do not cause any problem,
but they often form 10-50 microns sized agglomerates. These large agglomerates can
be a cause of concern in a casting. This is because TiB2 particles are relatively much
harder than aluminium. They can be detrimental to critical applications such as thin
sheet, micron foils, fine wire and applications requiring special surface finishes [72].
According to [57], TiB2 is considered to be a type of inclusions, whereas TiAl3, which
is soluble under normally encountered foundry conditions, is not.
1.3.3 The Assessment of Inclusion
The assessment of metal cleanliness is an integral part of any casting operation. The
optimum technique would provide the assessment of three inclusion parameters: size,
distribution, and composition. Several techniques of melt cleanliness are available to
assess the cleanliness of molten aluminum alloys [57,73,74]. Among the commonly
used techniques, the following can be mentioned:
The LiMCA (Liquid Metal Cleanliness Analyzer) technique [75,76] is an on-line melt
cleanliness determination, utilizing electrical resistivity to detect inclusions in the 20300 micron range, depending on the orifice size of the sampling tube used. This
technique is based on Electrical Sensing Zone (ESZ), as shown in Fig 1.2. A constant
electrical current is applied between the electrodes and flows through the liquid metal
in the small orifice of the tube. The presence of an inclusion in the melt flowing
through the orifice increases the electrical resistance measured at the orifice. The
signals are amplified and their amplitude is measured digitally. The advantage of this
technique is that the impact of furnace preparation, alloying practice, feedstock mix,
settling time and similar parameters on melt cleanliness, is easily determined. The
technique characterizes the cleanliness of a melt at time intervals of the order of 1
minute, i.e. it delivers an immediate response. The disadvantage of this technique is
that it also detects gas bubbles from degassing as particles, which complicates the
analysis, and it cannot provide information on the chemistry, shape, or the physical
state of the inclusions. The small orifice (~ 300μm) may become plugged or blocked,
if the metal contains relatively large inclusions. The cost of the measurement
instrument is expensive.
Figure 1.2: Schematic of the LiMCA principle.
The PoDFA (Porous Disc Filtration Analysis) [59,77] is an off-line technique,
developed by Alcan International Limited, in which a fixed volume of metal is forced
through a filter, concentrating the inclusions on the surface of the filter, as is shown in
Fig 1.3.
Figure 1.3: Schematic of the PoDFA Analysis 57.
Subsequent metallographic evaluation is needed to measure the amount and type of
inclusions present. While it is inexpensive to collect samples, this method requires
time and expertise to acquire the measurements. It provides information about the type
of inclusions and their relative proportions. However, in the PoDFA technique, since
the assessment involves concentration of the inclusions, some information is lost, such
as the distribution of the inclusions. Since the concentration also includes some degree
of agglomeration, the information about size and shape might also be lost.
Furthermore, the inclusion size is limited by the filter-porous disk size, i.e. small
particles cannot be detected by the PoDFA method.
Ultrasonic test is a non-destructive technique, and it transmits an ultrasonic wave
through the metal sample. Inclusions are identified through sound damping [78].
Considerable work has been carried out, using ultrasound, to detect inclusions in
molten aluminium alloys [78,79]. The disadvantage of this technique is that it remains
expensive and can only detect inclusions greater than 100μm.
Other less common and less widely used methods to measure the inclusion content of
metals are neutron activation analysis (used to measure oxygen concentration), arcspark emissions spectrometry [80], and selective dissolution techniques etc. [81].
However, these methods are also relatively expensive tedious and time-consuming.
The acid dissolution test, that is, extracting out the non-metallic inclusion by the
complete dissolution of the aluminium matrix in mixed acids and the collection of the
insoluble [82,83]. This test is limited by the difficulty of dissolving anything more
than a relatively small sample, and the fact that some of the compounds and
contaminants are also soluble in the solutions used.
When it comes to the assessments of inclusions in the solidified product, a small
number of inclusion assessment methods are known, and there is a lack of literature
data. A typical method used is the light microscopy examination. Obtaining results by
this method is, however, time-consuming, and at the time that the results are received,
the metal produced may already have left the house. Due to the low inclusion content,
and the high tendency of inclusion to segregate into a nonuniform distribution, the
direct microscopic analysis of the surface of metal samples is not practical, since a
very large area would have to be observed for quantitative assessment. In the present
thesis a simpler test for inclusion assessment in solidified Al alloy castings has
therefore been developed, using a NaOH macroetching [84,85]. One of the reasons for
the use of the macroetching technique is related to the random nature of the occurrence
of inclusions in Al alloys. Another reason could be attributed to the difficulties with
the detection and the quantitative assessment of inclusions.
1.4 The Scope of the Thesis
This study has been performed to provide an understating of the solidification
behaviour and the microstructure formation generated during the solidification of DC
casting of some industrially important wrought aluminium alloys. The intentions are
alloys development and to increase our understanding of the effects of the casting and
solidification parameters, the amount of certain alloying addition, and the dimensions
of casting products on the castability and resulting microstructures of the 3000 and
6000 series alloys in particular. An additional aim of the project was to simulate the
conditions (e.g. cooling rate) in the large DC casting process, using directional
solidification experiments with a Bridgman-type furnace and a DC simulator.
The project results will provide general guidelines to improve the quality and the
performance of the final products, to reduce casting difficulties and defects, to reduce
scrap, and to contribute on optimizing DC process parameters, with regard to
castability and ingot dimensions.
The thesis starts with an analysis of the thermal field and the solidification process in
samples during growth in a Bridgman furnace, in order to provide information about
the growth conditions, such as the temperature gradient, and the position and velocity
of the solid/liquid interface during directional growth [Suppl. 1]. The thesis also covers
the study of the solidification behaviour of aluminium wrought alloys and the effect of
excess additions of certain alloying elements on the microstructure formation [Suppl.
2-5]. The equilibrium solidification behaviour of the heat exchanger alloys of the
AA3003 type containing an excess of Zn and Cu were studied by a series of thermal
analyses (DTA, and DSC). The results of these experiments are presented in [Suppl.
4]. The last part of the thesis deals with quality issues, namely the location and content
of inclusions in DC-cast billets [Suppl. 6], and an application of the studied materials
to automotive production [Suppl. 7].
1.5 References
1. Stobart, Patrick D. Centenary of the Hall & Heroult Processes, 1886-1986. London: International
Primary Aluminium Institute, 138 pp, 1986.
2. Altenpohl, Dietrich G. Aluminum: Technology, Applications, and Environment: A profile of a
Modern Metal. Book: Aluminium from Within. 6th Edition. Publ: Warrendale PA 15086, USA;
Minerals, Metals & Materials Society (TMS); Nov.1997. Ed: J.G.Kaufman, S.K.Das.
3. Glenn Switkes. Foiling the aluminum indstry. A toolkit for communitles, Activists,Consumers,
workers. Published by International Rivers Network, August 2005.
4. Grandfield, J.F. and McGlade, P.T. DC casting of aluminium: process behaviour and technology.
Materials Forum, Australia. Vol. 20, pp. 29-51. 1996.
5. Grandfield, J.F. DC casting of aluminium: a short review of process development. Fifth Australian
Asian Pacific Conference. Aluminium Cast House Technology: Theory & Practice; Queensland;
Australia; June 1997. Pp. 231-243.
6. King, F. Aluminium and its Alloys. Ellis Horwood, Market Cross House, Cooper Street,
Chichester, West Sussex PO19 1EB, UK, pp. 313.1987
7. Roth, W. Deutsches Reichs Patent Nr. 974203. 1936
8. W. T. Ennor, US patent 2301027, Nov. 3, 1942.
9. Faunce, J P; Wagstaff, F E; Shaw, H. New Casting Method for Improving Billet Quality. Light
Metals, Los Angeles, Calif; U.S.A; 27 Feb.-1 Mar. 1984. pp. 1145-1158.
10. Jacoby, J. E. Direct chill casting defects. Fifth Australian Asian Pacific Conference. Aluminium
Cast House Technology: Theory & Practice; Queensland; Australia; June 1997. pp. 245-251.
11. Suyitno, Eskin D.G., Savran V.I., and Katgerman. L Effects of Alloy Composition and Casting
Speed on Structure Formation and Hot Tearing during Direct-Chill Casting of Al-Cu Alloys.
Metallurgical and Materials Transactions A. Vol. 35A, no. 11, pp. 3551-3561A. Nov. 2004
12. Eskin D G, Savran V. I, and Katgerman L. Effects of Melt Temperature and Casting Speed on the
Structure and Defect Formation during Direct-Chill Casting of an Al-Cu Alloy. Metallurgical and
Materials Transactions A. Vol. 36A, no. 9, pp. 2561A. Sept. 2005.
13. Eskin, D G; Zuidema, J; Savran, V I; Katgerman, L. Structure formation and macrosegregation
under different process conditions during DC casting. Materials Science and Engineering A. Vol.
384, no. 1-2, pp. 232-244. Oct. 2004.
14. Eskin D.G., Zuidema (Jr) J, Savran V.I. and Katgerman L. Effect of Process Parameters on
Structure Formation during Direct-Chill Casting of an Al-Cu Alloy. Materials Forum. Vol. 28, pp.
487-493. 2004.
15. Buxmann, K and Gold, E. Solidification Conditions and Microstructure in Continuously Cast
Aluminum. J. Met. Vol. 34, no. 4, pp. 28-34. Apr. 1982.
16. Anyalebechi, P N. Effects of alloying elements and solidification conditions on secondary dendrite
arm spacing in aluminum alloys. EPD Congress 2004 as held at the 2004 TMS Annual Meeting;
Charlotte, NC; USA; Mar. 2004. pp 217-233.
17. Wan, J. Lu, H-M, Chang, K-M, and Harris, J. As-Cast Mechanical Properties of High Strength
Aluminum Alloy. Light Metals; San Antonio, TX; USA; Feb. 1998. pp. 1065-1070.
18. Bunn, A. M, Schumacher, P, Kearns M. A, Boothroyd, C. B, and Greer, A. L. Grain refinement by
Al-Ti-B alloys in aluminium melts: a study of the mechanisms of poisoning by zirconium. Materials
Science and Technology (UK). Vol. 15, no. 10, pp. 1115-1123. Oct. 1999
19. Cooper, P and Barber, A. Review of the latest developments and best use of grain refiners. 2nd
International Melt Quality Workshop, Prague, Czech Republic, 15-17th October 2003.
20. Kashyap, K T; Chandrashekar, T. Effects and mechanisms of grain refinement in aluminium
alloys. Bulletin of Materials Science (India). Vol. 24, no. 4, pp. 345-353. Aug. 200.
21. Quested, T.E. (2004): Ph.D. Thesis. “Solidification of Inoculated Aluminium Alloys”. Department
of materials science and Metallurgy, University of Cambridge, Cambridge UK.
22. Kearns, M.A, Thistlethwaite S.R. and Cooper P. S. Recent Advances in understanding the
mechanism of aluminium Grain refinement by TiBAl master alloys. Light Metals; Anaheim,
California; USA; Feb. 1996. pp. 713-720.
23. Pearson, J and Cooper, P.S. A review of the basics of grain refining. Sixth Australian Asian Pacific
Conference on Aluminum Cast House Technology; Sydney; Australia; July 1999. pp. 109-118.
24. Easton M, StJohn D. (1999a): Metall Mat Trans A 1999; 30A: 1613.
25. Easton M, StJohn D. (1999b): Metall Mat Trans A 1999; 30A: 1625.
26. Mohanty P.S, Gruzleski, J.E. Mechanism of grain refinement in aluminum. Acta Metall. Mater.
Vol. 43, no. 5, pp. 2001-2012. May 1995.
27. Schumacher, P., Greer, A. L. (1994a). Heterogeneously nucleated alpha -Al in amorphous
aluminium alloys. Materials Science and Engineering A (Switzerland). Vol. A178, no. 1-2, pp.
309-313. Apr. 1994
28. Schumacher, P., Greer, A. L. (1994b): Mater. Sci. Eng. A, 182, 1335.
29. Schumacher, P., Greer, A. L., Worth, J., Evans, P. V., Kearns, M. A., Fisher, P., Green, A. H.
New studies of nucleation mechanisms in aluminium alloys: implications for grain refinement
practice. Materials Science and Technology (UK). Vol. 14, no. 5, pp. 394-404. May 1998.
30. Johnsson, M and Backerud, L. The influence of composition on equiaxed crystal growth
mechanisms and grain size in aluminum alloys. Zeitschrift fur Metallkunde (Germany). Vol. 87,
no. 3, pp. 216-220. Mar. 1996.
31. Meredith, M.W, Worthm J, Brown, J.M. and Hamerton, R. G. Understanding DC-cast
microstructures using laboratory-based solidification techniques. Light Metals 2003 as held at the
132nd TMS Annual Meeting; San Diego, CA; USA; Mar. 2003. pp. 1111-1118.
32. Yang X., Pekgüleryüz M.Ö., and Bouchard M. Modelling of Solidification of AL Alloys in a
Laboratory Scale DC-Simulator. Light Metals 2000 as held at the 129th TMS Annual Meeting;
Nashville, TN; USA; Mar. 2000. pp. 591-596.
33. Heiberg G. Gandin Ch.-A., Goerner H., and Arnberg L. Metallurgical and Materials Transactions
A. Vol. 35A, no. 9, pp. 2981-2991A. Sept. 2004
34. Maggs, S. J., Cochrane, R. F., Flood, S. C., Evans, P. V. The effect of trace element of
intermetallic phase selection in simulated dc castings. Light Metals 1995, Las Vegas NV, USA;
Las Vegas NV; USA; Feb. 1995. pp. 1039-1047.
35. Maggs, S. J. (1996): Ph.D. Thesis, University of Leeds, UK.
36. Meredith, M. W., Greer, A. L., Evans, P. V., Hamerton, R. G. Light Metals; San Diego, CA; USA;
28 Feb.-4 Mar. 1999. pp. 811-817.
37. Meredith, M. W. (1999): Ph.D. Thesis, University of Cambridge, UK.
38. Quested, T.E., and Greer, A. L. Modelling of microstructural evolution in Bridgman specimens.
Solidification of Aluminum Alloys as held at the 2004 TMS Annual Meeting; Charlotte, NC;
USA; Mar. 2004. pp. 43-52.
39. Greer, A.L., Quested, T.E., and Spalding, J.E. (2002): Modelling of grain refinement in directional
solidification. Light Metals 2002 as held at the 131st TMS Annual Meeting; Seattle, WA; USA;
Feb. 2002. pp. 687-694.
40. Speyer, R.F. Thermal Analysis of Materials. Marcel Dekker, Inc. (USA), pp. 297, 1994.
41. Tamminen, J. “Thermal Analysis for Investigation of Solidification Mechanisms in Metals and
Alloys”, Ph.D. Thesis, Univ. Stockholm, Sweden, 1988.
42. Starink M.J. Analysis of aluminium based alloys by calorimetry: quantitative analysis of reactions
and reaction kinetics. International Materials Reviews. Vol. 49, no. 3-4, pp. 191-226. June-Aug.
2004
43. Stefanescu, Doru M; Ruxanda, Roxana. Solidification Structures of Aluminum Alloys. Materials
Park, OH: ASM International, 2004. Pp. 107-115.
44. Polmear, I.J., Light Alloys- metallurgy of the light metals. London, 1989: Edward Arnold.
45. Altenpohl, D. Aluminum Viewed from Within. Aluminium-Verlag, 1982. ISBN 3-87017-138-3.
46. Karlsson, A. Braze Clad aluminium materials for automotive heat exchangers some recent
developments. INCAL '98: International Conference on Aluminium; New Delhi; India; Feb. 1998.
pp. 189-199.
47. Gray, Alan. The growth of aluminium in automotive heat exchangers. Aluminium. Vol. 81, no. 3,
pp. 197-201. Mar. 2005
48. Takemoto,T; Okamoto, I. Investigation on the Corrosion of Brazed Aluminium. Kei Kinzoku
Yosetsu. J. Light Met. Weld. Constr. Vol. 23, no. 1, pp. 2-13. Jan. 1985.
49. Itagaki, T; Toma, K. Corrosion Resistance of Al--Mn--Cu Alloys. 76th Conference of the Japan
Institute of Light Metals; Osaka; Japan; May 1989. pp. 125-126.
50. Davis, J.R. Corrosion of Aluminum and Aluminum Alloys. ASM International, Member/Customer
Service Center, Materials Park, OH 44073-0002, USA, 1999. pp 313.
51. Hirth, S.M., Marshall, G J, Court, S A, Lioyd, D J. Effects of Si on the aging behaviour and
formability of aluminium alloys based on AA6016. Materials Science and Engineering A
(Switzerland). Vol. 319-321, pp. 452-456. Dec. 2001
52. Shoji, M; Yamauchi, S. Influence of Titanium Additive on the Corrosion Resistance of Al--Mn-Mg--Si--Cu Based Alloys. 76th Conference of the Japan Institute of Light Metals; Osaka; Japan;
May 1989. pp. 127-128.
53. Lessiter, M J, and Rasmussen W.M. To pour or not to pour - The dilemma of assessing your
aluminum melt's cleanliness. Modern casting, February 1996. 45-48.
54. Lessiter, M J. Understanding Inclusions in Aluminum Castings. Modern Casting (USA). Vol. 83,
no. 1, pp. 29-31. Jan. 1993
55. Eklund, J.E. On the effect of impurities on the solidification and mechanical behaviour of primary
and secondary commercial purity aluminium and aluminium alloys-PhD thesis. Helsinki
University of technology, Espoo, Finland. 1991. 135P.
56. Eckert, C E; Miller, R E; Apelian, D; Mutharasan, R. Molten Aluminum Filtration: Fundamentals
and Models. Light Metals 1984; Los Angeles, Calif ; U.S.A ; 27 Feb.-1 Mar. 1984. pp. 1281-1304.
57. Doutre, D, Gariepy B, Martin J. P and Dube, G. Aluminum Cleanliness Monitoring: Methods and
Applications in Process Development and Quality Control. Light Metals 1985; New York, New
York; USA; Feb. 1985. pp. 1179-1195.
58. Simensen, C.J and Berg, G. A Survey of Inclusions In Aluminum. Aluminium. Vol. 56, no. 5, pp.
335-338. May 1980
59. Apelian, D. How clean is the metal you cast? The issue of assessment; a status report. 3rd
International Conference on Molten Aluminum Processing, Florida; USA; Nov. 1992. pp. 1-15.
60. Eckert, C.E. The origin and identification of inclusions in foundry alloys. 3rd International
Conference on Molten Aluminum Processing; Orlando, Florida; USA; Nov. 1992. pp. 17-50.
61. Apelian, D and Mutharasan, R. Filtration: a melt Refining Method. J. Met. Vol. 32, no. 9, pp. 1418. Sept. 1980
62. Gallo, R. Correlating Inclusion Sizes with Aluminum Casting. Modern Casting. Vol. 94, no. 6, pp.
27-30. June 2004
63. Gallo, R. Development, Evaluation, and Application of Granular and Powder Fluxes in Transfer
Ladles, Crucible, and Reverberatory Furnaces. 6th International AFS Conference on Molten
Aluminium Processing; Lake Buena Vista, FL; USA; Nov. 2001. 15 pp. 2002.
64. Martins, L. C. B., and Sigworth, G. K. Inclusion Removal by Flotation and Stirring. 2nd
International Conference on Molten Aluminum Processing; Florida; Nov. 1989. pp. 16.1-16.28.
65. Groteke, D.E. Influence of SNIF Treatment on Characteristics of Aluminum Foundry Alloys.
Transactions of the American Foundrymen's Society, Vol. 93; Pittsburgh, Pennsylvania; USA; 29
Apr.-3 May 1985. pp. 953-960.
66. Schmahl, J.R. Aubrey L. S. and Martins, L.C.B. Recent Advances in Molten Aluminum Filtration
Technology. 4th International Conference on Molten Aluminium Processing; Orlando, FL; USA;
Nov. 1995. pp. 71-102.
67. Groteke, D.E. The Reduction of Inclusions in Aluminum by Filtration. Mod. Cast. Vol. 73, no. 4,
pp. 25-27. Apr. 1983
68. Neff, D. V. Continuous, "re-usable" filtraton systems for aluminum foundries. 4th International
Conference on Molten Aluminium Processing; Orlando, FL; USA; Nov. 1995. pp. 121-137.
69. Crepeau, P N. Molten aluminum contamination: gas, inclusions and dross. Modern Casting
(USA). Vol. 87, no. 7, pp. 39-41. July 1997
70. Campbell, J.C. Castings 2nd Edition (Book). Oxford, Butterworth Heinemann, 2003. 335 pp.
71. Fuoco, R., Correa, E.R., Bastos, M., and Escudero, L.S., “Characterization of Some Types of
Oxide Inclusions In Aluminum Alloy Castings”, AFS Transactions, v. 107, pp. 287-294 (1999).
72. Whitehead, A. J., Cooper, P. S., and McCarthy, R. W. An Evaluation of Metal Cleanliness and
Grain Refinement of 5182 Aluminum Alloy DC Cast Ingot Using Al-3%Ti-0.15%C and Al-3%Ti1%B Grain Refiners. Light Metals 1999; San Diego, CA; USA; 28 Feb.-4 Mar. 1999. pp. 763-772.
73. Dasgupta, S., Apelian, D. Interaction of initial melt cleanliness, casting process and product
quality: cleanliness requirements fit for a specific use. 5th International AFS Conference on
Molten Aluminium Processing; Orlando, FL; USA; Nov. 1998. pp. 234-258.
74. Simard, A., Prouix, J; Paguin, D; Samuel, F.H. Capability study of the Improved PREFILFootprinter to measure liquid Aluminium Cleanliness and Comparison with Different Techniques.
6th International AFS Conference on Molten Aluminium Processing; Lake Buena Vista, FL; USA;
Nov. 2001. 20 pp. 2002
75. Guthrie, R.I.L. and D.A. Doutre, D.A. On-Line Measurements of Inclusions in Liquid Melts.
Refining and Alloying of Liquid Aluminium and Ferro-Alloys. The Norwegian Institute of
Technology; Trondheim; Norway; Aug. 1985. pp. 145-163.
76. Ware, TN; Cooksey, M; Couper, M J. Measurement of the performance of in-line processes using
LiMCA. Seventh Australian Asian Pacific Conference on Aluminum Cast House Technology;
Hobart; Australia; Sept. 2001. pp. 45-54.
77. Utigard, T.A. and Sommerville, I.D. CLEANLINESS OF ALUMINUM AND STEEL: A
COMPARISON OF ASSESSMENT METHODS. Light Metals 2005 as held at the 134th TMS
Annual Meeting; San Francisco, CA; USA; Feb. 2005. pp. 951-986.
78. Gallo, R; Mountford, H; Sommerville, I. Ultrasound for On-Line Inclusion Detection in Molten
Aluminum Alloys: Technology Assessment. AFS International Conference on Structural Aluminum
Castings: Recent Advancements for the Design and Manufacturing of Reliable Structural
Aluminum Castings; Orlando, FL; USA; Nov. 2003. pp. 179-194.
79. Mansfield, T L.Ultrasonic Technology for Measuring Molten Aluminum Quality. Light Metals;
Dallas; Tex; Feb. 1982. pp. 969-981.
80. Semensen, C J. Gas-Chromatographic Analysis of Carbides in Aluminium and Magnesium.
Fresenius Z. anal. Chem. Vol. 292, no. 3, pp. 207-212. 1978.
81. Simensen, C.J. Sampling and Analysis of Impurities in Aluminium. Refining and Alloying of Liquid
Aluminium and Ferro-Alloys. The Norwegian Institute of Technology; Trondheim; Norway; Aug.
1985. pp. 165-171.
82. Siemensen, C.J., and Strand, G. Analysis of Inclusions in Aluminium by Dissolution of the Samples
in Hydrochloric/Nitric Acid. Fresenius Z. Anal. Chem. Vol. 308, no. 1, pp. 11-16.
83. LaVelle, D. L. Determining Abrasive Inclusions in Aluminium Alloy Castings. Foundry/ October
1969. pp 154-157.
84. Vander Voort. Metallography: Principles and Practice. ISBN: 0-87170-672-5
85. Petzow G., Metallographic Etching, ASM (American Society for Metals), 1978, P. 39.
Chapter 2
2
Experimental methods
2.1 Materials
The materials used in the thesis were aluminium alloys of mainly AA3003 and 6xxx
series, provided by Kubikenborg Aluminium AB. The standard chemical compositions
of these types of alloys are given in Table 2.1.
The material was DC-cast as billets or slabs with different dimensions. Samples for
remelting investigations were taken from a region of 10 mm away from the edge of the
ingots in order to get representative compositions.
Table 2.1: Chemical compositions (wt%) of aluminium alloys, 3xxx and 6xxx series.
Element
Mn
Fe
Si
Mg
Cu
Zn
Al
3xxx
0.9-1.5
<0.7
<0.6
0.037
<0.2
0.1
Balanced
6xxx
0-0.8
0-0.8
0.4-1.3
0.3-1.2
Balanced
2.2 Experimental Techniques
The following laboratory techniques were used in this thesis to simulate industrial
scale DC casting:
2.2.1 Directional Solidification
Directional solidification methods are widely used techniques to simulate and study
fundamental issues associated with the solidification during industrial DC casting.
These methods include the Bridgman technique and the DC simulator. The DC
simulator provides a range of solidification conditions over the length of the samples,
which replicate the conditions present in full-scale DC casting. Bridgman
solidification allows a series of samples to be cast that can cover the complete range of
solidification conditions found in full-scale DC casting.
Measurements of temperature were made within these techniques by inserting
thermocouples at representative positions; the recorded temperature data obtained
were then used to determine the thermal field during solidification and to validate
modelling of the process.
2.2.1.1 Bridgman technique
The Bridgman directional solidification apparatus used for this study is illustrated
schematically in Fig. 2.1. It consists of three zones, a hot zone with a potassium heat
pipe, an adiabatic zone about 0.2 m long, and a cold zone with water filled heat pipe.
The furnace is equipped with transition device and quenching facility.
Figure 2.1: Schematic of the Bridgman experimental equipment.
The Bridgman samples were prepared from the DC cast ingots by remelting in alumina
crucibles at 800 oC, with subsequent casting in copper moulds of 5 mm diameter. The
cast rods of a total length of 115 mm were then transferred into alumina tubes of 5 mm
inner diameters. The specimens were melted in the Bridgman furnace in an argon
atmosphere at 800oC by drawing the furnace 7 cm downwards with 2mm/min melting
speed. The specimens were held at that position for one hour to ensure stabilization of
the solidification front. Growth was then driven by moving up the furnace at different
constant pulling rates. The furnace motion was controlled by motor and gearbox, and
the movement was monitored by means of a pointer, which was attached to the furnace
and moves along a vertical scale. Applied pulling speeds were in the range 1 to
50mm/min. After a certain pulling distance of the furnace ~ 30-55 mm, the
solidification was interrupt as the sample was quenched by dropping it downward into
water, located under the furnace.
The measurements of temperature profiles during Bridgman solidification were carried
out in some experiments by inserting thermocouples within the samples of the
composition 3003 aluminium alloys. A K-type thermocouple wire, 0.5 mm diameter
was inserted into a thin alumina protection tube, 0.8 mm inner diameter, and 1.6 mm
outer diameter. These were placed within the Bridgman samples in a hole drilled
inside the sample close to the wall of the alumina tube surrounding the sample. The
position of the thermocouple tip was; either 50 mm, 80 mm or 100 mm from the
bottom of the sample as seen in Fig. 2.2.
For each thermocouple position, the sample was placed in the Bridgman furnace in a
standard position. By further analysis of the recorded temperature data it was possible
to calculate the temperature gradient, position and velocity of the solidification front
and cooling rate during the growth. The temperature gradient inside the sample was
about 4.7 °C/mm at the solidification front. The cooling rate for the 50 mm min–1
pulling rate was in the range of 0.1-1.2 °C/sec, and a corresponding analysis for the
lowest pull rate, 1 mm/min, gave cooling rates in the range between 0.05-0.1 °C/sec.
Figure 2.2: Schematic of locations of the thermocouple.
In the Bridgman solidification, during pulling at the high rate, 50 mm/min for 55 mm,
the growing interface did not follow the furnace pull rate. The solid/liquid front had
only moved about 10 mm. The actual growth rates were calculated from the
measurements of the thermal field data, and if the average growth rates are estimated
as the ratio of solidification distance over solidification time, the resulting average
growth rates become similar.
2.2.1.2 DC simulator
A method that can simulate industrial DC casting processes is the chilled casting test
apparatus, also knows as DC-simulator, in which, the solidification occurs in a
direction approximately perpendicular to a chill surface. The cooling rates close to the
chill surface are approximately similar to those in the shell zone of a DC-cast ingot.
Figure 2.3: Schematic illustration DC simulator apparatus.
The experimental set up used for the DC simulator is illustrated in Fig 2.3. It consists
of a copper block (100mm x 100mm x 30mm) isolated from the sides by ceramic
insulator plates lining of 20 mm thickness, and a casting cup (120mm x 80mm x
30mm), made of steel that also served as a riser and was positioned directly on top of
the copper mould. The metal was poured into alumina tubes of 12 mm outside
diameter and 8 mm inner diameter, which were positioned in the copper mould to
shape the melt and avoid transverse heat transfer to the copper mould. Directional
solidification was achieved by cooling the bottom surface of the copper plate.
The thermal field was determined during solidification by temperature measurements
using K-type thermocouples placed at heights, 5 mm and 25 mm above the chill
surface. The cooling rate at 5 mm from the chill surface was 2 oC/s, the solidification
rate 1 mm/s, and the temperature gradient 2 oC/mm.
The conditions, (e.g. velocity, cooling rate and thermal gradient) were changed a long
the length of the sample throughout the solidification process in a way similar to the
variation through the thickness of a commercial ingot. For example, the cooling rate
measured closer to the bottom of the sample is approximately similar to the cooling
conditions present at the surface of a commercial ingot, whereas slightly further into
the sample, the cooling rate was measured to be much lower. Since the conditions are
changing constantly through the ingot, so do aspects of the microstructure.
2.2.2 Thermal Analysis Techniques
Differential thermal analysis (DTA) and differential scanning calorimetry (DSC)
experiments are widely utilized techniques to simply and accurately measure the
temperatures of phase reactions in materials, as well as to study small as-cast samples
with controlled cooling rate. These calorimetric techniques have been proved to be a
useful (complementary) tool in microstructure characterization and their results are
used to obtain information regarding the solidification behavior and precipitation
reactions occurring in alloys on heating and cooling.
2.2.2.1 Differential thermal analysis (DTA)
Differential thermal analysis (DTA) is a technique, which measures the difference in
temperature between the investigated sample and a reference material (a thermally
inert material- graphite), which are subject to the same heating or cooling as a function
of time. When a sample undergoes a phase transformation, it will either absorb,
endothermic process, or release heat, exothermic process. The integrated temperature
difference as a function of time is a measure of the amount of heat that is absorbed or
released during the phase transformation. A schematic diagram of a DTA set-up is
shown in Fig 2.4.
Figure 2.4: Schematic of DTA set-up.
DTA specimens of a cylindrical shape of about 13 mm length and 8 mm diameter were
heated to 800oC under an argon atmosphere. The experimental data were collected
during melting and solidification at heating and cooling rates of 10 and 4 oC/min
(0.167 and 0.067K/s). DTA curves were obtained using a sample holder and with two
thermocouples (S-type). The sample temperature was measured at the centre of the
sample and a reference temperature was measured in a graphite sleeve.
During solidification some samples were quenched in water at temperatures that have
been chosen according to earlier recorded DTA cooling curves to investigate the start
of the solidification, their reactions and formed microstructures.
2.2.2.2 Differential Scanning Calorimetry (DSC)
DSC is a technique determining the variation in the heat flow exchanged between the
sample and the surrounding as a function of temperature or time. At heating or cooling
transformations taking place in the material are accompanied by heat exchange. DSC
determines the temperatures of these transformations and quantifies their heat
exchanges, i.e. enthalpy changes in the alloys. Fig. 2.5 shows a schematic of a
Differential Scanning Calorimetry instrument of heat flux type DSC. It consists of a
reference pan and a sample pan contained in a constant temperature enclosure.
Figure 2.5: A heat flux type DSC (left) Heat source (right) Sample carrier.
The measurements were performed using a NETZSCH model DSC 404 with discshaped samples having a weight of approximately 70 mg encapsulated in alumina pan.
Sapphire (62 mg) was used as a reference. During the measurements, the samples were
protected with flowing argon. Temperature scans were made in the range 50-700oC
with constant heating and cooling rates of 20, 10, 4, and 1 K/min. The amount of
energy, which has to be supplied to or withdrawn from the sample to maintain zero
temperature differential between the sample and the reference is the experimental
parameter displayed as the ordinate of thermal analysis curve.
2.2.3 Testing
2.2.3.1 Mechanical testing
Hardness testing has always been important in the metalworking industry. Measuring
hardness is a relatively quick, cheap and easy way to assess the strength of a material
without the need to prepare tensile test samples. For example, it may be a convenient
way of investigating the progress of precipitation hardening.
The micro hardness measurements were performed using a Vickers micro hardness
test apparatus under a known load. Pressing a point with a diamond tip on the sample
surface with a known force, the size of the diagonal is taken as a measure of the
hardness. In the thesis the hardness testing was carried out to evaluate the mechanical
properties. In the Bridgman samples, the hardness has been tested with a small load
(50 N), in the centre of the dendrites. In the DC simulator samples, the overall
structure was tested using a load of 300 N. An average of minimum 10 hardness
indentations on each sample surface was measured.
Compressions strength tests have been also performed at elevated temperatures to
assess the strength of a material. Samples prepared for hot compression were cut 15
mm height and machined to completely flat edges. They were taken 5, 20 and 35 mm
from the chill surface to measure the effect of different cooling rates.
Isothermal compression tests were performed using a MTS load frame. The testing
was conducted at temperatures 400, 450 and 500°C and at strain rates of 0.01 and 1/s.
A Lenton resistance furnace, regulated with temperature controllers, was used for
heating two water-cooled parallel platens in which two thermocouples (k type) were
installed to monitor the temperature of the specimen. Oil based graphite was used as
the lubricant between the specimen and the platens to prevent frictional effects during
the test. When the platens had settled at the desired test temperature, a specimen was
clamped between them and held at this temperature for three minutes before
compression. Tests were run using constant velocity control profiles. Load and stroke
were acquired during the compression test by a data acquisition system and later
converted to true stress-true strain curves.
2.2.3.2 Macroscopic Etching Test
A macroetching technique has been used in this thesis to study inclusions in solidified
aluminium billets. It was employed to satisfy easy handling, easy sampling and quick
evaluation of inclusions in a large area at low cost. The principle idea is to employ
visual inspection of macroetch surface of a test piece; a sample of the ingot cross
section, and count the number of inclusion appearing as etch pits on the surface
regardless of its size, kind and type. The technique can be a useful tool in the
evaluation of the cleanliness of the melt, and the inclusions content and distribution
along aluminium cast billets.
A disk, cut in transverse direction of a DC-cast Al billet, is grounded and macroetched
by immersing in hot NaOH solution. The conditions are 10-20 g Sodium hydroxide in
100 ml water, at 60 to 70°C, for about 5 to 15 min. After etching, every kind of
inclusions such as, Al2O3, MgO, MgAlO4, TiB2, etc, generate holes, i.e. etch pits in the
surface that can be seen by the naked eye. The surface condition for parts to be
macroetched is not critical. A finely polished surface is not required; a flat surface
obtained by a turning machine is adequate. However, the surface should be free of
grease or oil to permit even etching.
To easy characterize the inclusion distribution on billet slices, mapping of the etched
surfaces were made by using a COPY machine, and the copies were then scanned to a
computer. The number of inclusions that become visible as black dots were then
counted using an image analysis program (KAPPA software). Fig. 2.6 shows an
example of pictures of the lateral distribution of pits on etched billet cross sections.
2.5cm from billet bottom block
56cm from bottom block
2.5cm from billet top
Figure 2.6: Image of the etched billet cross sections.
In Fig. 2.7a-b it can be seen by optical micrographs how an inclusion developes,
during subsequent etch steps, into a large etch pit, visible by the naked eye. It is the
etching of a TiB2 clusters that is followed.
No etch
10sec etch
90sec etch
15min etch
30sec etch
Figure 2.7a: Subsequent etch steps of a TiB2 cluster.
No etch
10sec etch
30sec etch
90sec etch
5.5min etch
15min etch
Figure 2.7b: Subsequent etch steps of an oxide+TiB2 clusters.
2.2.4 Metallographic investigations
Microstructure characterization of the investigated aluminium alloys was carried out
by optical and electron microscopy. Measurements of microstructure features, such as
secondary dendrite arm spacings (SDAS), grain size and the chemical compositions of
the secondary phases were carried out. The investigations were made both on as-cast
samples of direct chill (DC) castings and on specimens solidified in a Bridgman
furnace or under controlled solidification conditions.
The solidified samples were cut, grinded on 320 SiC-grit papers and polished with
suitable diamond suspension agents, 9 μm, on a hard polishing cloth followed by
polishing on a soft polishing cloth with 3 µm and 1 µm diamond paste and finally a 35 minutes treatment with an OPS-solution. Then, in order to reveal the structure,
etching was carried out with a reagent using, HF solution, (0.5 ml HF in 100 ml water)
for 20 seconds. After etching, the dendrite structure was revealed. For easy grain
measurements, the samples were anodized with a platinum anode in a medium of 24%
Ethanol, 74%H2O, 1%HF, 1%HBF4 and direct current of 50 V for 90 seconds.
The directionally solidified samples were prepared for optical microscopy analysis
using longitudinal sections. For the DC cast ingots, samples for microstructural
examinations of the size 5 x 15mm were cut at appropriate distances from the surface
toward the centre of the ingot.
An optical microscope with attached camera, and software programs called Image–
Pro, and/or KAPPA Image were used in the metallographic analysis of the
microstructure features. The secondary dendrite arm spacings were measured by
averaging the distance between adjacent side branches of a primary dendrite. The
measurements of grain sizes were performed in light optical microscopy using
polarized light, as the number of grains on the circumference of a circle.
In addition to the light optical microscope investigations, a scanning electron
microscope (SEM) equipped with an energy dispersive x-ray (EDX) detector was used
to provide information about the chemical compositions of the precipitated secondary
phases and inclusions types in the investigated alloys. In the presented work
transmission electron microscopy (TEM) was also used primarily to determine
inclusion particles, which were difficult to detect in SEM due to their small size.
Because in many instances no well-developed dendritic structure i.e. dendritic crystals
growth with a clear side-branching structure, could be revealed; a technique of
measuring the distances between secondary phases and liquid films was applied to
obtain numerical values of the coarseness of the structures. This was done by placing
lines randomly in the structures and measuring the average free distances along the
lines.
Chapter 3
3
Results and discussion
3.1 Results of thermal conditions in the Bridgman furnace
Temperature measurements within a sample of aluminium AA 3003 alloys in a
Bridgman type furnace during growth (and no-growth conditions) were performed in
Suppl. 1. The technique was employed to determine the evolution of temperature field
and also to provide information about the growth conditions such as, the temperature
gradient, position and velocity of the s/l interface during the directional growth. Thus
it makes it possible to understand the performance of the used furnace and to adjust the
furnace parameters to simulate the range of solidification conditions that may be found
in industrial-scale DC casting.
Table 3.1: Compositions of grain refined Al 3003 alloys (wt.%).
Alloy
3003
Mn
1.14
Fe
0.50
Si
0.135
Mg
0.002
Ti
0.018
Cu
0.126
Zn
0.008
Temperature, oC
The results of temperature measurements at three thermocouple positions 50, 80 and
100 mm from the bottom of a sample during solidification by furnace pulling with
different speeds 1, 10 and 50 mm/min are shown in Figure 3.1. The temperature was
recorded at regular intervals approximately 5mm apart during the furnace motion. By
further analysis of the recorded temperature data and mathematical treatment, it was
possible to calculate the temperature gradient, position and velocity of the
solidification front and cooling rate during the growth.
850
100mm
750
80mm
650
50mm
1mm/min
10mm/min
50mm/min
1mm/min
10mm/min
50mm/min
1mm/min
10mm/min
50mm/min
550
450
350
250
0
10
20
30
40
50
60
70
Furnace traverse distance, mm
Figure 3.1: Temperatures recorded during furnace pulling as a function of pulled distance.
In the mathematical treatment, quadratic functions of the type shown in Equation 3.1
were fitted to the temperature data selected at the three thermocouple positions at
different time interval and thus, the temperatures could be obtained as a function of
time for all positions in a sample.
T (t ) = C1 X 2 + C2 X + C3
(3.1)
Cooling curves could thus be obtained for differed pulling rates. Fig. 3.2 shows the
calculated cooling curves at a pulling rate of 50 mm/min as an example.
Temperature (oC)
850
800
750
700
650
600
550
0
0,2
0,4
0,6
0,8
1
1,2
1,4
Time (min)
45mm
70mm
100mm
50mm
80mm
110mm
60mm
90mm
115mm
Figure 3.2: Calculated cooling curves at 50mm/min pulling rate.
By using the temperature data, the positions of an isotherm can be plotted as a function
of time. This has been done in Fig 3.3, where both the solidus and the liquids
isotherms have been plotted for two different pulling rates.
1,8
1mm/min VL=0.99mm/min
VS=0.97mm/min
90
1,6
1,4
1,2
80
1
70
0,8
60
0,6
IL
IS
VL
VS
50
40
0
10
20
30
40
50
0,4
Front velocity mm/min
Interface Position, mm
100
0,2
0
60
Furnace pulling time, min
Figure 3.3a: Interface position and velocity versus time.
50mm/min
Interface position, mm
58
56
54
25
VL=8.14mm/min
VS=7.82mm/min
20
15
52
50
10
48
46
44
42
IL
5
IS
VL
0
VS
40
0
0,5
1
Front velocity, mm/min
60
-5
1,5
Furnace pulling time, min
Figure 3.3b: Interface position and velocity versus time.
The slopes of the linear fit of s/l location curves give values for average growth rates
as listed in the Fig 3.3. The derivative of each points in the location curves give values
of the instantaneous growth rates as a function of time for the studied pulling
velocities, Fig. 3.3.
Table 3.2, shows the operating conditions of some of the directional solidification
experiments for the furnace test. It is clear that there is a deviation between growth
distances and pulled distances, i.e. the average growth rate was not equal to the
furnace-pulling rate at higher pull rates.
Table 3.2: Operating conditions for directional solidified experiments.
Furnace
Pull
distance
at
quenching
(mm)
55
70
40
40
10
30
Furnace
Pulling rate
[mm/min]
50
50
8
5
1
1
Grown
distance
in
solidified
sample
(mm)
8
11
25
33
7
31
Growth
Time
[min]
1.1
1.4
5
8
10
30
Calculated
average
growth rates
(mm/min)
7
8
5
4
0.7
1.0
Typically the growth rate starts with a transient, which is followed by a period of
constant rate, but at the end, the growth rate increase again.
The temperature gradient at the s/l front can also be obtained from the temperature
data as the derivative of the quadratics function at the liquidus and solidus
temperature. The values of the gradients at the melt interface for the three studied
furnace speed, are plotted in Fig. 3.4. It can be observed that the gradient decrease
with distance and time during the growth. The maximum gradient of approximately
6oC/mm is observed at the beginning of the solidification.
Temperature gradients,
o
C/mm
6,5
1mm/min
10mm/min
50mm/min
6
5,5
5
4,5
4
3,5
3
2,5
40
50
60
70
80
90
100
Distance from the bottom of the sample, mm
Figure 3.4: Calculated temperature gradient as a function of distance.
A one dimensional heat flow model (Equ. 3.2) could account for first order effect with
regard to the found growth rate variations along the samples. The temperature gradient
in the liquid at the interface, GL get the value 4.6 oC/cm at the low pull rate. This is close to
4.8, which was obtained in Fig 3.4 in the stable gradient zone, i.e. 65-80mm from the bottom
of the sample. The analytical heat flow modelling can therefore successfully be applied to
experiments that reach steady state, but for higher growth rates the boundary conditions for
equation (3.2) are not met.
Vo
yl
T −T
V
GL = hotVo cold . o e α L
⎛ yl ⎞ α L
2⎜⎜ e α L −1⎟⎟
⎝
⎠
(3.2)
Microstructural analysis from Bridgman experiments showed clearly how the
solidification parameters affect the resultant microstructure. Equiaxed microstructure
observed in samples solidified with high growth rate and mixed columnar-equiaxed
microstructures were observed at the lowest solidification rate, Fig. 3.5.
Figure 3.5: The shape of the solidification front for high and low growth rates respectively.
The directional solidified microstructures were compared with the DC-cast
microstructure. It was observed that the majority of solidification in Bridgman samples
adopts a similar morphology to that in DC cast, Fig 3.6.
Figure 3.6: Microstructures in DC cast ingots and in a directional solidification samples faraway
from solidification front respectively.
3.1.1 Conclusions
The objectives of this work were to investigate whether the thermal conditions in a
Bridgman furnace, were comparable to the conditions during DC-casting. It was found
that the used Bridgman furnace quite well could be adapted to DC-casting simulation
of certain parts of the ingots.
In Bridgman solidification, steady-state solidification can be achieved over a larger
proportion of the length of the sample by using longer samples; in the current set-up,
this was limited by the sample manufacturing and by the size of the furnace.
3.2 Results of different Mg and Si contents on solidification of 6XXX alloys
In the case of Al 6xxx alloys, it is Mg and Si that form the strengthening precipitates.
Alloy with higher contents of Si than Mg gives a good strength and extrudability.
However the casthouse experience is that such alloys are more sensitive to surface
defects, and thus more difficult to cast than alloys with balance content of Si and Mg.
The influence of the balance between Si and Mg concentrations on the microstructures
of some important extrusion alloys, namely 6063, 6005A and 6082 were investigated
for the aim to increase the knowledge of their solidification behaviour, and to study the
effect of addition of excess Si on the microstructure of AA 6063 alloy. The
investigations were made for the compositions given in Table 3.3.
Table 3.3: Compositions of the investigated alloys (wt.%).
Alloy Type
6063
6063+Si
6082
6005A
Mg
0.4
0.42
0.64
0.55
Si
0.43
0.62
1.01
0.70
Fe
0.17
0.17
0.20
0.21
Mn
0.02
0.02
0.48
0.13
Ti
0.012
0.010
0.024
0.011
Cu
0.001
0.003
0.010
0.105
The as-cast microstructures of the three DC cast ingots 6063, 6005A, 6082 were
analysed by optical microscope, and the grain sizes were measured along the diameter
of billet slices. No significant difference in gain structures was observed between the
three alloys as can bee seen in Fig. 3.7.
Grain size [Microns]
200
150
6063
6005A
6082
100
50
0
40
80
120
Radial distance [mm]
Figure 3.7: The grain size as a function of the radius of the three studied billets.
The microstructures of as-cast samples of the investigated 6005A and 6082 alloys at
the surface zone are shown in Fig 3.8a. The surface structures show significant
differences in the two alloys surface region. The 6005A alloy has a wider enriched
segregation region with large precipitations in grain boundaries, while the secondary
phases at the 6082 surface seem to concentrate strongly to a more narrower segregate
layer and have the appearance of plate like precipitations.
The investigated alloys were also studied in experiments performed by the Bridgman
technique, which simulate the growth in the large commercial ingots. Samples from
the three ingots were directionally solidified at pulling velocities between 5–50 mm
min–1, (0.83 10-4 - 8.3 10-4 m/s), which give cooling rate in the order of 0.25-1.0oC/s.
The temperature gradients inside the sample were varied, and were 25oC/cm and
35oC/cm for the two-furnace temperatures 800 and 1000oC respectively.
40μm
40μm
6005A
6082
Figure 3.8a: Surface structure of 6005A and 6082 ingots.
Results of the Bridgman experiments showed that, increased Si contents result in
larger grain size. Fig. 3.9 shows the grain size (GS) in quantitative form as a function
of Si content, and Fig. 3.10, shows that the directionally solidified microstructures
close to the quenched front were clearly influenced by the Si content.
Grain Size (GS), [Micron]
200
GS, 50mm/min
GS, 5mm/min
150
100
50
0,4
0,5
0,6
0,7
0,8
0,9
1
1,1
Si content in Alloy [wt%]
Figure 3.9: Grain size as a function of Si contents.
200μm
0.43% Si
0.62% Si
1.01% Si
Figure 3.10: Directionally solidified structures close to the quenched front, pulling rate
50mm/min.
Figure 3.11. Shows the structure coarsening from quenched front to start of growth for
two directionally solidified samples of 6063 alloy with different Si contents. It can
clearly be observed in Figure 3.11 that Si additions promote development of secondary
dendrite arms and result in a coarser structure with a larger grain sizes, but with shorter
distances between secondary phases. The final structure at room temperature is much
coarser than the structure formed at solidification, implying that the transformations
during cooling are very important for the final structure and the materials properties.
200μm
0.43% Si
0.62% Si
Figure 3.11: Structure coarsening from quenched front (top) to start of growth (bottom) for two
directional solidified samples at 50 mm/min pulling speed.
3.2.1 Conclusion
Differences have been noted in the investigated three ingots; especially the ingots
surface structures showed significant variations, which probably influences the surface
defect formation. The mechanisms for that is not yet clear, but the inverse segregation
to the surface also plays an important roll giving significantly higher alloy contents in
this region.
3.3 Results of additions of Zn and Si on the solidification of 3003 alloys
Zinc is often chosen as an alloying element because; it is highly soluble and can be
dissolved in solid state in aluminium to high concentrations. Thus, the strength can be
improved by solution hardening. It is expected that application of Zn additions to
current aluminium heat exchanger materials, 3003 type alloys can provide a significant
improvement in the corrosion performance of heat exchangers. Si is used for strength
improvement as it contributes to both particles and solid solution strengthening.
The effect of the additions of Zn and Si on the microstructure formation of 3003 Al
alloys was investigated by Bridgman directional solidification technique with
quenching were applied. The chemical compositions of the base alloys are as tabulated
in Table 3.4.
Table 3.4: Composition (wt.%) of AA3003 alloys prior to additions.
Alloy
AA3003
A
B
Dimension
mm x mm
1440x330
1440x600
Mn
Fe
Si
Mg
Ti
Cu
Zn
Zr
1.23
1.160
0.500
0.510
0.145
0.489
0.037
0.009
0.019
0.018
0.146
0.125
0.021
1.514
0.004
0.002
Alloy (A) was used in experiments with Zn additions and alloy (B) was used in the
experiments with Si additions. The different levels of Zn additions were 1.5, 2.5, and
5.2 wt.% and the Si addition to alloy (B) was 2 wt.%.
Significant effects on the structures of the alloys formed during directional
solidification experiments were found as a result of varying amount of alloy additions.
For the zinc additions the microstructure became finer i.e. smaller SDAS and grain
size as the Zn content increased to 2.5%, but at 5.2 wt% a coarser structure was
formed. In Fig. 3.12 the quenched solidification fronts in the Bridgman samples are
shown. In the 2.5% alloy, a larger number of less developed dendrites appear giving a
finer grain size. Also between the dendrite arms more secondary phases were precipitated in this alloy, which hinders the coarsening after the solidification resulting in a
final structure being finer both with regard to dendrite arm spacing and grain size.
0.1mm
0.02% Zn
1.4% Zn
2.5% Zn
5.2% Zn
Figure 3.12: Microstructures of quenched interfaces in Bridgman samples with different Zn level.
Additions of Si, which have a significantly lower solubility than Zn, quickly give a
finer structure, larger amounts of secondary phase precipitations and a longer
solidification interval. Fig. 3.13 shows the refining effect of Si additions.
400μm
0.47% Si
1.9% Si
Figure 3.13: Microstructures along the Bridgman solidified samples from quenched front (top) to
start of growth (bottom) with two different Si contents.
For Zn additions in the range of 0-5 wt.%, the hardness measured by a Vickers micro
hardness test apparatus, showed clear effect, Fig. 3.14. In the Bridgman samples,
hardness values from the centre of dendrite arms increased with increased Zn content,
as Zn goes into solid solution. In the DC simulator samples, also containing 0.35-0.9
wt% Cu, the hardness values from the overall structure, show a clear peak at 2.5 wt.%
Zn as a result of both largest fraction of θ-phase and finest structure observed at this
composition.
110
Bridgman sample
DC simulator-5mm
DC simulator-25mm
Hardness (HV)
100
90
80
70
60
50
40
30
0
1
2
3
4
5
6
wt% Zn
Figure 3.14: Hardness of alloys with different Zn contents. Bridgman samples measured in centre
of dendrites, and DC simulator samples measured in overall structure by larger indentations.
3.3.1 Conclusion
Solidification studies of Zn additions to Al 3003 alloys showed that the most important
effect on microstructure was noticed at 2.5 wt.% Zn where the structure was fine and
the hardness had a maximum. Si addition to a level of about 2% gave a finer structure,
and a relatively large fraction of eutectic structure, but also a long solidification
interval, which limits the use of Si as a strengthen element to about 0.6%.
3.4 Results of DTA and DSC studies of 3003 alloys containing Zn and Cu
The purpose of this paper was to give a better understanding of the influence of
different Zn and Cu additions on the solidification behaviour, phase precipitation and
resulting microstructure in Al 3003 alloys. The cooling curves were recorded during
the solidification in DTA and DSC experiments. The alloys compositions were Al
3003 alloy with Zn additions in range 0-5wt.% and 0.15-1.0wt% Cu, Table 3.5.
DTA curves obtained for 3003 alloys containing about 0.15 wt.% Cu and different Zn
additions during solidification at a constant cooling rate of 4 K/min (0.067oC/s) are
shown in Fig. 3.15. For the three lower Zn concentrations, i.e. 0.02, 1.5 and 2.5 wt%,
only one second peak appears in the curves, giving the temperatures for the start of
secondary phase precipitation. The values of these temperatures for each alloy are
listed in Fig. 3.15. For 5.2% Zn alloy the peak is not clearly visible, and only a slight
flattening of the curve, which indicates a small amount of secondary phase formation,
can be seen. The liquidus, the solidus and the temperatures of the secondary phase
precipitation are decreasing with increasing percentage of Zn contents in the alloys.
The values of these characteristics are listed in Table 3.5. Similar effects, i.e. melting
point depression were also observed for the alloys with different Cu contents.
ΔT=T s-T r /oC
16
14
0.02 Zn
12
1.5Zn
10
2.5Zn
8
5.2 Zn
Tsecondary
628
precipitation
639
645
6
T0.02=645
T1.5=642
T2.5=640
T5.2=634
4
To=651oC
2
0
-2580
590
600
610
620
630
640
650
660
670
o
T sample / C
Figure 3.15: DTA cooling curves of 3003 alloys for different Zn additions cooled at a rate of
4oC/min. The open circles on the DTA thermogram indicate where samples were quenched.
DSC cooling curves recorded at the rate of 4oC/min are presented in Fig. 3.16, which
shows the effect of varying Zn and Cu on the investigated 3003 alloys.
0,5
DSC /mW/mg
0
600
-0,5
-1
-1,5
610
620
630
640
650
660
670
0.02%Zn+0.15Cu
2.5Zn+0.15Cu
2.5Zn+0.94Cu
2.3%Zn+0.35%Cu
5.2%Zn+0.15Cu
-2
Temperature /oC
Figure 3.16: DSC cooling curves of 3003 alloys with different Zn and Cu content at a cooling rate
of 4 K/min.
In Table 3.5, it can clearly be observed that the increased Zn content and the presence
of excess Cu have a significant effect on the melting point depression, on the
solidification ranges, and temperatures of the secondary phase precipitation. Especially
at the highest alloy contents, i.e. in the samples containing 5.2Zn+0.15Cu and
2.5Zn+1.0Cu, the melting point depressions are obvious.
Table 3.5: Data from DTA and DSC experiments at rate of 4 K/min
Alloys
name
0.02Zn
1.5Zn
2.5Zn
5.2Zn
2.5Zn0.35Cu
2.5Zn1.0Cu
TLiquid (OC)
DTA
DSC
651
657
649
648
654
645
650
648
653
644
650
Tp-precipitation (OC)
DTA
DSC
645
642
640
634
640
636
640
637
635
637
635
TSolid (OC)
DTA
DSC
641
638
635
625
634
624
630.5
626
621
625
622
T-range (OC)
DTA
DSC
10
11
13
20
14
20
27
28
29
27
28
The results of the microstructure investigations of the as-quenched DTA samples with
different Zn are shown in Fig 3.17. It can be seen how the different crystallization
morphologies change as the Zn and Cu contents increase. At 2.5% Zn, a globular
structure, i.e. less developed dendrites were formed. At 5.2% Zn a coarser dendritic
structure with larger solidification intervals was formed. The analysis of the alloys
with different Cu contents, Fig. 3.18, shows that an increase of the amount of Cu gives
a more developed dendritic growth morphology.
800μm
(a)
(b)
(c)
Figure 3.17: Optical microstructures of the quenched DTA samples (a) 1.5Zn+0.15Cu
(Tquench=645), (b) 2.5Zn+0.15Cu (Tquench=639oC) and (c) 5.2Zn+0.15Cu (Tquench=628oC).
800μm
(a)
(b)
(c)
Figure 3.18: Optical microstructures of the quenched DTA samples contain 2.5%Zn and different
Cu. (a) 2.5Zn+0.15Cu (Tquench=639oC) (b) 2.5Zn+0.35Cu (Tquench=642oC) and (c) 2.5Zn+1.0Cu
(Tquench=635oC)
Although only one second peak can be seen on the DTA cooling curves, two
secondary phases were observed in the investigated microstructures. Type 1 was
identified by EDX-analysis in a SEM to be Al6(MnFe) and type 2 was α–
Al15(Fe,Mn)3Si2, which is formed both from the liquid and as a transformation of the
first phase. Fig. 3.19a shows the appearance of these phases in the base alloy and Fig.
3.19b shows a partly transformed particle.
(a)
(b)
Figure 3.19: SEM images of (a) the precipitated phases in the base alloy and (b) shows the
transformation where α–Al(Fe,Mn)-Si is brighter than the Al6(Fe,Mn) phase.
EDX-analysis also indicated how the Zn and Cu entered into, and stabilized the
different phases. Generally only Zn goes into the type 1-phase as Al6(Fe,Mn,Zn) but
both Zn and Cu can be found in the type 2-phase as Al15(Fe,Mn,Zn,Cu)3Si2, i.e. at 45% Zn, the Al6(Fe,Mn,Zn) phase is dominating, and Si is mainly found in the Al
matrix.
EDX analysis of structures formed by DC simulator test at higher cooling rate than the
DTA samples indicated a similar effect occurred but even stronger. Figure 3.20 shows
clearly how the amount of secondary phases increase with the Zn content until 2.5%,
but upon further Zn additions it decreases again. Also the dendrite arm spacing has a
minimum at 2.5%, and the SEM pictures reveal a segregation pattern in the 1.5 and 4%
Zn structures, but not in the 2.5% alloy.
0.02%Zn
2.5%Zn
1.5%Zn
4%Zn
Figure 3.20: Backscatter SEM microstructures showing the effect of different Zn contents
EDX analysis of the DC simulator structures also reveal that an eutectic phase in the
2.5% Zn and 1.0% Cu alloy, has a composition corresponding to Al2(Cu,Zn,Si)
resembling the appearance of the θ-phase, Al2Cu, in the Al-Cu system. In the 4% Zn
alloy the dominating phase is Al6(Fe,Mn,Zn) with only small amounts of the light
Al2(Cu,Zn,Si) phase. This is in agreement with the DTA results, where Al6(Fe,Mn,Zn)
is strongly dominating at higher Zn contents.
3.4.1 Conclusions
The DTA and DSC analysis showed clearly the differences caused by the additions of
Zn and Cu to the balance 3003 alloys. The results of microstructural investigation
showed that at 2.5% Zn, globular structure was formed, whereas the alloy with 5.2%
Zn has a coarser structure and relatively large solidification range.
In all the DTA and DSC samples, where the cooling rates are low, and the
solidification processed close to equilibrium conditions; only two secondary phases
were observed in the alloys. Neither Zn nor Cu did form own intermetallic phases but
were found in the solid solution and in precipitated phases as Al6(Fe,Mn,Zn) and
Al15(Fe,Mn,Zn,Cu)3Si2. However ingot solidification occurs at a higher cooling rate,
and therefore at increased Zn contents in the presence of Cu, the Al2(Cu,Zn,Si) start to
form at 1.5% and has a maximum at 2.5%, while at higher Zn concentration
Al6(Fe,Mn,Zn) is the dominating phase.
3.5 Results of additions of excess titanium
Titanium is usually added to many aluminium alloys because it is a potent grain
refiner, i.e. gives smaller grain size, and at higher additions can possibly improve the
corrosion resistance. However the casthouse experience shows that if Ti is added as an
alloying element and not only as a grain refiner to the level of the peritectic point and
above 0.15-0.20 wt.%, makes it more difficult to cast the alloy without cracks.
The effect of different Ti concentrations was studied by both characterizing the grain
structure over the thickness of DC cast rolling ingots, and by studying the
solidification microstructures of Bridgman directionally solidified samples. The
investigated alloy compositions, as shown in Table 3.6, were Al 3003 alloys produced
by DC casting. The grain refiner was Al-5%Ti-1%B.
Table 3.6: Compositions of the investigated alloys.
Alloy
Dimension
mm x mm
1440X400
1440x460
1440x460
1440x600
1440x600
1440x600
Mn
Fe
Si
Mg
Ti
Cu
Zn
Zr
0.76
0.843
0.931
1.14
1.18
1.55
0.198
0.216
0.209
0.5
0.53
0.258
0.142
0.139
0.139
0.135
0.147
0.748
0.212
0.002
0.083
0.002
0.009
0.013
0.148
0.155
0.163
0.018
0.029
0.013
0.283
0.305
0.276
0.126
0.115
0.042
0.010
0.006
0.011
0.008
0.041
1.429
0.143
Grain size [Microns]
Figure 3.21, shows the grain sizes for varying Ti levels, measured from the surface to
the centre of several DC cast ingots. It can be seen that, the high Ti alloys have a
coarser structures than the alloys with the normal Ti concentrations for grain
refinement, i.e. <0.02 %Ti. Fig. 3.21 also shows the poisoning effect of Zr, i.e. the
grain refining mechanism cease to work.
600
500
+Zr
400
300
200
100
0
0
100
200
300
Distance from ingot surface [mm]
Ti=0.148wt%
Ti=0.155wt%
Ti=0.163wt%
Ti=0.018wt%
Ti=0.029wt%
Ti=0.013%+Zr
Figure 3.21: The as-cast grain size as a function of distance from the ingot surface for different Ti
contents and for an alloy with Zr.
Experiments were done by Bridgman directional solidification with high and normal
Ti contents. It was found that the high Ti alloys had a finer grain size and stronger ten-
dency to equiaxed growth, Figure 3.22, which is opposite to the results of the DC cast
ingots where the high Ti alloy showed a coarser grain structure.
The results were explained by an analysis of the dissolution time of the Al3Ti particles,
which showed that the dissolution time increased from about 60 to 200 seconds when
the Ti content increased from 0.015 to 0.15%. If the Al3Ti particles do not totally
dissolve before the melt starts to solidify, a few large Al3Ti particles will nucleate the
aluminum grains instead of the numerous boride particles, which in this case are
smaller and not activated as nucleation points.
Figure 3.22: Microstructures of quenched solidification fronts for low Ti and high Ti respectively,
solidified at 50mm/min pulling rate. Solidification direction upwards with quenched liquid on top.
This conclusion was supported by an analysis of the bulk ingot structure. Some Al3Ti
particles could be revealed in the centre of aluminium dendrites, as shown in Figure
3.23. The undissolved large particles are therefore likely the reason to the coarser
structure and the higher cracking sensitivity in the high Ti alloys. The reason why the
structures in the Bridgman experiments were finer at the higher Ti content was that in
these cases the time for dissolution of Al3Ti particles was long enough and therefore
the high Ti level could increase the growth restriction and give a finer structure.
100μm
0.155% Ti
100μm
0.163% Ti
Figure 3.23: Al3Ti at the centers of grains in the microstructure of DC-cast ingots.
3.5.1 Conclusion
Comparison of alloys of different Ti contents showed that; addition of titanium to a
level > 0.02% gives a coarser grain structure than alloys with a lower Ti content. This
is probably a result of that the Al3Ti particles in grain refiner master alloys are not
totally dissolved during the casting process implying nucleation on larger but few
particle. The larger grains sizes in the high Ti alloys are most likely the results of high
crack sensitivity during DC casting of these alloys.
3.6 Results of assessment of inclusions in aluminium billets
The knowledge of the amounts, distribution and locations of inclusions along DC cast
billets length and cross section is of large industrial importance. A good assessment
method could result in reduced scrap, i.e. an improvement of the yield by decreasing
the cutting off length at the ingot ends and scalping depth at the ingot surface. Due to
the low inclusion content, and the high tendency of inclusion to segregate into a
nonuniform distribution, the direct microscopic analysis of small metallographic
samples is not practical, since a very large area would have to be observed for
quantitative assessment.
Macroetching technique based on NaOH acid solution has been developed to study
inclusions in solidified aluminium billets. It was employed to satisfy easy handling,
easy sampling and quick evaluation of inclusions in a large area at low cost. Four
billets of two different types of alloys, AA 6063 and AA 6082 were analysed. Two of
the billets namely, A1 and A2 were of the same charge but from different positions of
the casting table, while the billets named B, and C were from different charges.
The technique is that, a disk of the billets cross-section is immersed in hot NaOH
solution. After the deep etching, inclusions generate holes/etch pits on the surface. The
results of subsequent etch steps in Fig. 3.24 show how a TiB2 inclusion develops into a
visible hole in the surface.
no etch
10sec etch
30sec etch
90sec etch
5.5min etch
15min etch
Figure 3.24: Subsequent deep etching of an oxide+TiB2 cluster inclusion.
An example of a macroscopic view of the inclusion distributions on deep etched crosssections of billets is presented in Fig 3.25. The photographs show a non-uniform
lateral distribution of holes (black dots), which correspond to inclusions. A high
inclusion content in the billets surface regions can generally be seen compared to
interior. It can also be seen that the first billet disk at position 2.5 cm from bottom
block, shows a large number of very small holes, taking a shape of circular zone close
to the mould wall.
12cm
2.5cm
60cm
788cm
Figure 3.25: Photographs of the etched billet cross sections, showing the lateral distribution of
etch pits along the length of a billet (samples position from the bottom block).
The quantity of inclusions on the etched surfaces of different ingots was measured
regardless of the pits type and kind by an image analysis program. The holes were
classified based on their maximum length. Three size ranges were used and results are
presented in Fig. 3.26. It shows the size distribution changes along those billets.
Higher densities of etch pits are revealed at both ends, i.e. at the bottom and top parts,
compared to the main part of the billet in between. At the bottom samples, i.e. 0-3 cm
from the bottom block, the highest densities of inclusions are located. At 5-12 cm the
relative amount of large etch pits increased compared to the rest of the billet.
100
A1: 0.2<L<0.3
A2: 0.2<L<0.3
B: 0.2<L<0.3
Numbers of holes (length)
90
80
A1: 0.3<L<0.4
A2: 0.3<L<0.4
B: 0.3<L<0.4
A1: L>0.4
A2: L>0.4
B: L>0.4
70
60
50
40
30
20
10
A1
A2
B
0
4.5-5
11-12.
56-57
786-790
788-792
Sample position from the bottom end along the billets (cm)
Figure 3.26: The variations of etch pits counted in transverse sections along three billets of two
different charges.
Optical microscopic charactarization were performed on selected pieces 20mm x 10
mm, which were cut from the studied macroscopic etched sections. In the slices from
2.5 cm from the bottom block, the majority of discovered imputities were aluminium
oxides, (Al2O3) and borides cluster (TiB2). At ~56 cm form the bottom block,
magnesium oxides were found in additions to Al2O3, and TiB2. Fig 3.27 shows typical
appearnaces of the different inclusion types in these two positions.
(a)
(b)
Figure 3.27a: Optical micrograph showing inclusions of (a) Oxide particles and (b) oxides films in
sample position 2.5 cm from bottom block of the billet A1.
(a)
(b)
(c)
Figure 3.27b: Optical micrograph showing inclusions (a) spinal MgAl2O4 (b) MgO and (c) TiB2
cluster in sample position at 56 cm from bottom block of the billet A1.
3.6.1 Conclusions
The results of the studied billets revealed similar distribution tendency of the etch pits
along the billets length. A trend in decreasing number of etches pits from the bottom
block toward the top with an increase again at the very top of the billets. Also the
presence of a ring shaped distribution of a large number of small defects in the
beginning of the casting, i.e. just above the bottom block could be revealed. These
higher densities of pits in the beginning of the casting are related to the disruption of
the oxides films due to the strong turbulence during pouring the liquid aluminium to
the mould. The slight increase in the inclusion frequency at the very top part of the
billet can be explained by the effect that at the end of the casting process the melt flow
is disturbed as the melt level in the launder system start to decay. The ring shaped
small inclusions are identified as TiB2 particles and associated with the grain refinment
proceedure. During casting practice excess of grain refiners are often added in the start
of the casting. The general high etch pits density closer to the billets surface might be
couple to the melt flow conditions and the high solidification in this resgion.
4
General Conclusions of Thesis
The solidification process and the resulting microstructure are very important for the
mechanical properties of the cast materials. By increasing the basic knowledge of the
crystallisation process it will be possible to improve the properties and tailor the
material for different applications. The properties of aluminium alloys are influenced
by a number of microstructural features, which are controlled by chemistry and
processing.
The present study was carried out to investigate the role of some major operating
parameters applied in aluminum foundries and that of minor alloying element
additions on microstructure, properties, and quality in 3003 and 6xxx wrought
aluminium alloys, using laboratory casting techniques.
From an analysis of the results, the following general conclusions may be drawn.
• Addition of alloying elements such as Zn, Si, Cu, and Ti can be used to improve
properties and tailor the material for different applications, if the amounts are
optimized.
• Another way of improving the properties is by employing optimum additions of
grain refiners.
• Inclusions are a problem, and better techniques to detect them and their
distribution, will improve the possibilities to limit the problem.