Effects of intensive forced melt convection on the mechanical properties of Fecontaining Al-Si based alloys
X. Fang, G. Shao*, Y.Q. Liu, Z. Fan
BCAST (Brunel Centre for Advanced Solidification Technology)
Brunel University, Uxbridge, Middlesex, UB8 3PH, UK
* Email address: [email protected] Tel.: +44 1895 266967; Fax +44 1895 269758.
Abstract
Fe is the most common and usually detrimental impurity in cast Al-alloys. As the equilibrium solid
solubility of Fe in the Al solid solution (α-Al) is rather low, Fe exists in Al-alloys in the form of Febearing intermetallic compounds. In commercial Al-Si based cast alloys, these compounds are often
of the morphologies as either long needles or large plates, which drastically deteriorate the ductility
of the alloys. The Fe impurity in Al-alloys results mainly from the use of steel tools and scrap
materials, which is an inherent engineering problem for industrial use and recycling of Al-alloys.
This work aims to understand the effects of intensive forced convection on the mechanical
properties of Fe-containing Al-Si based cast alloys. Varied Fe levels were introduced into
commonly used commercial Al-Si based cast alloys with different Mn contents. The role of melt
convection was investigated through comparative study of the microstructure-properties
relationship of alloy test bars that were processed via rheo-diecasting (RDC) and conventional high
pressure die-casting (HPDC) respectively, with the former involving intensively shearing the melt
in a slurry maker before diecasting. CALPHAD modelling of the thermodynamic properties of the
multi-component alloys is carried out to study the role of alloying on the formation of different
primary Fe-bearing compounds, the cubic αAlFeMnSi with a compact morphology, or the βAlFeSi
with a needle- or plate-shaped morphology. The results show that the Fe tolerance of Al-Si based
cast alloys can be significantly extended by combining intensive forced melt convection with
adequate Mn/Fe ratios.
Keywords: Al-alloy, rheo-diecasting, CALPHAD, Fe-bearing compounds, mechanical properties
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1. Introduction
Commercial Al-alloys always contain Fe, often as undesirable impurity and occasionally as a useful
minor alloying element. As a minor alloying element, Fe has been used occasionally in Al-Cu-Ni
alloys to enhance high temperature mechanical properties, in Al-Mg alloys to reduce coarsening, in
Al-Fe-Ni alloys to enhance corrosion resistance in steam [1], and in high pressure die-cast (HPDC)
Al-alloys to facilitate ejection and to help die-release [2, 3]. In industrial practice, the impurity Fe
pickup results mainly from the use of steel tools during melting and casting, and the use of scrap
materials. For the vast majority of Al-alloys such as Al-Si based cast alloys, the presence of Fe is
detrimental to mechanical properties [1, 2, 4], and Fe pickup is thus a major problem for industrial
use and recycling of Al-alloys. Extensive efforts have to be made to keep the Fe levels as low as
economically possible.
The detrimental effects of impurity Fe to Al-alloys are due to its low equilibrium solubility in the
α-Al solid solution phase (< 0.04 wt.%) [1], and the associated strong tendency to form various
low-symmetry Fe-bearing aluminides. When these low symmetry compounds crystallise, in
particular, as primary phases during solidification, they are prone to grow into long needles/plates
that are extremely detrimental to both strength and ductility. Ternary Al-Fe-Si phases that are of
particular importance to Al-Si based cast alloys are the hexagonal αAlFeSi (Al8Fe2Si, Pearson
symbol hP246, space group P63/mmc) and the monoclinic βAlFeSi (Al5FeSi, monoclinic lattice)
[5]. The latter is particularly harmful to Al-Si based cast alloys, especially when it is formed as
primary phase during solidification. Therefore, the conventional metallurgical solution to eliminate
the detrimental effect of Fe has been largely focused on chemical means, either to limit the
maximum Fe content to avoid the formation of primary Al-Fe-Si compounds, or to modify the
crystal structures of Al-Fe-Si compounds into higher-symmetry lattice types.
The ternary eutectic reaction L → ( Al ) + Si + βAlFeSi occurs at 576 °C with a liquid composition
of Al-12wt.%Si-0.7wt.%Fe [1,6], and the maximum Fe content to avoid the formation of primary
βAlFeSi is therefore rather low. Understandably, for commercial Al-Si based cast alloys, the Fe
content is limited to be below 0.7wt.% [6,7]. Another approach is to introduce Mn to Al-Si-Fe
alloys, so that a cubic ternary Al15Mn3Si2 (or αAlMnSi) compound is stabilised to consume Fe by
forming an equilibrium quaternary phase, Al15(Fe,Mn)3Si2 (to be referred as αAlFeMnSi in this
work: Pearson symbol cP138, Space group Pm3 ) [5]. Cr may also be added as an Fe corrector,
often together with Mn [1]. The cubic αAlFeMnSi will solidify into compact morphologies due to
its high-symmetry crystal structure [1, 8, 9].
The first study of the Al-Fe-Mn-Si system was dated back to 1942 by Philips and Varley, who
investigated quaternary Al-Fe-Mn-Si alloys over the range of 0 to 4 wt.% of Mn, Fe and Si [10]. In
this region, they did not detect any quaternary phases. Later, Phragmén found that the cubic
quaternary αAlFeMnSi phase was in equilibrium with the monoclinic βAlFeSi phase in the Al-FeMn-Si system [11]. They claimed that this quaternary phase is a continuous solid solution phase
between αAlFeSi and αAlMnSi. Barlock and Mondolfo pointed out that it was unlikely to form a
continuous solid solution phase between them, as these two phases are of different structures [12].
Barlock and Mondolfo established a quaternary section running through Al3Fe, Al6Mn and Si [12].
They concluded that the αAlFeMnSi phase forms via continuous substitution of Mn with Fe in the
αAlMnSi lattice, and the Mn contents in αAlFeSi and βAlFeSi are rather limited [12]. Zakharov et
al. constructed some vertical sections in the range of 10 to 14 wt.% Si, and 0 to 4 wt.% of Fe and
Mn [13]. They reported a stoichiometry of Al16(Fe,Mn)Si3 for the αAlFeMnSi phase and confirmed
low Mn solubility in the αAlFeSi and βAlFeSi phases [13]. They proposed that the solubility of Fe
in the αAlMnSi phase is low (up to 1.8 wt.%), and the equilibrium Mn solubility in the αAlFeMnSi
quaternary phase is as low as 4 wt.%. This contradicts previous results [12] in that Fe substitution of
Mn in the αAlMnSi phase is discontinuous, or phase separation occurs in this structure concerning
mixing of Fe on the Mn sublattice.
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More recently, Devignon et al. studied phase equilibria at the Al-rich corner of the Al-Fe-Mn-Si
system at 550 °C using micro-probe analysis under scanning electron microscope, and reported that
the equilibrium compositions of the quaternary αAlFeMnSi phase can contain 72.7 to 74.7 at.% Al,
9.1 to 13.7 at.% Fe, 2.8 to 5.7 at.% Mn, and 9.5 to 11.5 at.%Si [9]. This recent work is consistent
with the Al15(Fe,Mn)3Si2 stoichiometry, which is based on Fe substitution of the Mn sublattice sites
of the Al15Mn3Si2 lattice, leading to idealised Al and Si contents of 75 at.% and 10 at.%
respectively. Overall, the Mn solubility in the ternary αAlFeSi or βAlFeSi compounds is very
limited (below 0.5 at.%) [9,12,14], and for the quaternary αAlFeMnSi phase, Fe replaces the Mn
sublattice and there is limited mixing in the Si sublattice. The stoichiometry of the αAlFeMnSi
phase is thus best described by Al15(Fe,Mn)3Si2, with the crystal structure of the ternary αAlMnSi
(Al15Mn3Si2).
Although it has been realised since the 1960’s that the detrimental effect of Fe on mechanical
properties can be eliminated or corrected through Mn addition, limited work has been directed
towards a systematic study. While the work of Barlock and Mondolfo [12] has been used widely as
a guideline for the Mn effect, practically no commercial Al-alloys fall within the field of their
Al3Fe-Al6Mn-Si pseudo-section. Recently, Shabestari studied the effects of Mn/Fe ratios on
primary intermetallic phases in an Al-12.7wt.%Si alloy with Fe and Mn levels up to 1.2wt.% and
0.5wt.% respectively [8]. Preliminary thermodynamic modelling of the Al-Fe-Mn-Si system has
been carried out by Balitchev et al., who showed that the results of Zakharov et al. could be
described reasonably well by using the COST507 database [15] for light metals and treating the
αAlFeMnSi as a stoichiometric compound [14], though the stoichiometry of Al16Fe1.7Mn2.3Si3 that
they assigned to the αAlFeMnSi phase contained remarkably higher Mn (10at.%) than experimental
findings (2.8 to 5.7 at.%) in quaternary Al-Fe-Mn-Si alloys near the Al-Si edge [9].
It will lead to continual build-up of transition metals (Fe, Mn, Cr etc) hence increased quantities
of intermetallic phases to correct the detrimental effect of Fe through controlling alloy chemistry
alone, particularly when recycled materials are used. Therefore, for sustainable management of
materials recycling, addition of corrector elements (e.g. Mn, Cr) should be as low as practically
possible. It is therefore desirable that the detrimental effect of Fe be removed by combining alloy
chemistry with physical means, so that the morphologies of Fe-bearing phases can be modified with
minimal addition of corrector elements.
Recently, researchers at BCAST of the Brunel University have developed an innovative semisolid
processing technology, rheo-diecasting (RDC), for processing of Al-alloys [16]. It consists of two
basic functional units, a twin-screw slurry maker and a standard cold chamber high-pressure
diecasting (HPDC) machine. The twin-screw slurry maker has a pair of screws with specially
designed profiles to achieve intensive melt shearing at high shear rate and high intensity of
turbulence. The screws and barrel are made from special materials in order to prevent reaction with
molten Al. In comparison with conventional HPDC, RDC offers superior grain refinement,
remarkable structure homogeneity, significantly reduced porosity and segregation, and thus
improved mechanical properties [17-19]. Additionally, RDC is effective in modifying phase
morphologies such that solid solution phases can solidify into spherical or globular particles in the
sheared melt slurry, leading to microstructures containing dendrite-free primary solid solution
phases [18-20]. It would be desirable to explore the effects of melt shearing on the morphologies of
Fe-bearing intermetallic compounds.
The present study attempts to evaluate the combined effects of Mn and melt shearing on the
morphology and distribution of different Fe-bearing compounds in widely used commercial Al-Si
based cast alloys (LM24 and LM25). The role of alloy chemistry on the Fe-effect has been
investigated by CALPHAD (acronym of Calculation of Phase Diagrams) modelling of multicomponent Al-Cu-Fe-Mg-Mn-Si-Zn system. The resultant thermodynamic database was used to
guide the selection of RDC processing parameters, in order to carry out RDC with a semi-solid
slurry that contains a chosen volume fraction of primary solid phase/s, i.e. the fcc α-Al solid
solution phase and Fe-bearing intermetallic compounds, and to understand the role of alloying on
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combating the detrimental effect of Fe (A set of criteria for thermodynamic guidance of RDC was
presented in ref. [21]). The effects of RDC on the microstructure-properties relationship were
studied with respect to samples processed with conventional HPDC. The results show that it is
effective to modify the morphologies of Fe-bearing intermetallic phases by combining Mn alloying
addition with semi-solid melt shearing, which leads to evidently enhanced Fe tolerance and
improved mechanical properties.
2. Experimental
The alloys with different Fe contents were produced by melting appropriate quantities of
commercial LM24 and LM25 ingots and an Al-46.45 wt.% Fe master alloy. The small amounts of
the master alloy used in this work had little effect on alloy compositions other than the Fe levels.
The LM24 and LM25 based materials in this work correspond to two different levels of Mn
contents, i.e. 0.30 and 0.16 wt.%Mn for LM24 and LM25 based alloys respectively (Table 1). The
overall alloy compositions were analysed using optical mass spectroscopy and confirmed by area
energy dispersive X-ray (EDX) quantification.
Tensile test samples are produced by a 2800kN die-cast machine for both conventional HPDC and
RDC. For RDC in this work, alloy melts were introduced into the slurry maker at 780°C, where
they were sheared isothermally for 45 seconds at a constant shear rate of 890 s-1 with a shearing
temperature of 580 °C for LM24 and 590 °C for LM25 based alloys. The slurries were then die-cast
at a die temperature of 220 °C to produce standard tensile samples. The pouring melt temperature
for HPDC was 640 °C with the die temperature being the same as that for the RDC.
Microstructures of the samples were studied by optical microscopy and scanning electron
microscopy. The latter was performed together with EDX analysis using a field-emission electron
source operating at 15 keV. EDX quantification was performed on polished samples, and the
libraries of standard X-ray profiles for EDX were generated using pure elements. In situ
spectroscopy calibration was performed in each session of EDX quantification using pure copper.
Care was taken to guarantee adequate spatial resolutions for EDX quantification, such that reliable
quantification was only carried out for phases with sizes greater than 2µm.
The tensile test was conducted according to ASTM standard B557 using the ASTM standard
tensile samples of 6.4 mm diameter and 25mm gauge length [22]. All the tests are performed at
room temperature using an Instron 5569 Table Mounted Materials Testing System with a ±50kN
load cell and a strain rate of 1.0×10-3.
3. Results
3.1 Microstructures of the LM24 based alloys
3.1.1 Microstructures of the HPDC samples of the LM24 based alloys
Different types, and amounts, of primary intermetallic Fe-bearing phases were related to the levels
of Fe contents. Little intermetallic compound was present in the starting LM24 material that
contains 0.7wt.%Fe. The alloy with the Fe content of 0.97wt.%, LM24_0.97wt.% Fe, contained
only a small amount of spherical intermetallic compounds, which were identified by EDX
quantification as the αAlFeMnSi phase with compositions typical of the αAlFeMnSi phase (Table
2). On the other hand, alloys with higher Fe contents (1.34 to 2.04 wt.%) contained needle-shaped
βAlFeSi in addition to the αAlFeMnSi phase as primary intermetallic compounds, with the amount
of the former increasing with Fe content and that of the latter nearly constant. The primary
αAlFeMnSi particles exhibit a spherical morphology when they are small, and they may develop
into more complex morphologies (e.g. cross-like or star-like) with increasing sizes, which is
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consistent with the αAlFeMnSi morphologies reported in previous work [9, 12]. The βAlFeSi
compounds exist as large needles or plates. For example, an overall view for the coexistence of the
primary αAlFeMnSi and βAlFeSi compounds is shown in Fig. 1a. An enlarged view in the vicinity
of large primary βAlFeSi needles is shown in Fig. 1b.
The α-Al solid solution phase exhibited a typical dendritic morphology (e.g. Fig. 1b). There was
no significant change in the morphologies of α-Al and eutectic structures for alloys of different Felevels. The interdendritic regions are characterised with a eutectic-like microstructure (labelled as
“E” in Fig. 1b). The eutectic structures are made of typical Al/Si eutectic phases and transitionmetal-bearing intermetallic compounds. Intermetallic phases could be easily identified by EDX
quantification, due to their significantly different compositions.
The EDX results of compositions for different phases are listed in Table 2. The αAlFeMnSi and
the βAlFeSi exist mainly as primary compounds. The major difference in the compositions of the
αAlFeMnSi and βAlFeSi phases is in their Mn and Si contents (Table 2), making them readily
distinguishable by both appearances and morphologies. The low Mn content (about 0.6 at.%) in the
βAlFeSi phase is consistent with previous findings [9, 14]. In contrast to the monoclinic βAlFeSi
phase, the cubic αAlFeMnSi phase was found to contain some Cu and Cr as well. Cu was found to
exist mainly in the interdendritic θAl2Cu.
3.1.2 Microstructures of the RDC samples of the LM24 based alloys
The RDC microstructures consisted of primary solid phases that are formed in the slurry-maker
during shearing in the semisolid phase region, and secondary microstructures that are solidified in
the subsequent stages after the slurry is transferred to the shot chamber for die-casting [20]. In the
RDC samples, there were two types of morphologies for the α-Al solid solution (e.g. Fig. 2a).
Globular or spherical α-Al particles with sizes ranging from 20 to 60 µm in diameters were isolated
by regions of finer dendritic (radial sizes ~ 5 to 10 µm) α-Al and their associated interdendritic
eutectic microstructures (labelled as “E”). These two types of α-Al morphologies corresponded to
phases formed at the different stages of solidification, with the former occurred in the slurry maker
during melt shearing and the latter formed during the stage of subsequent die-casting. Following
convention of definitions of previous publications [20], we classify the former as primary α-Al and
the latter secondary α-Al. The effect of RDC on the primary α-Al solid solution phase was to
modify the dendritic α-Al morphology into globular primary particles.
Comparing RDC microstructures to that of their HPDC counterparts, it is seen that the effect of
RDC on the morphologies of primary Fe-bearing compounds was significant. In alloys containing
up to 1.34 wt.%Fe, the long needle-shaped βAlFeSi primary phase was largely eliminated from the
RDC microstructures. Even in the alloy containing up to 2.04wt.%Fe, most primary Fe-bearing
intermetallic compounds were in the form of globular αAlFeMnSi (Fig. 2), with rough branching or
star-like features of these in the HPDC samples being largely smoothed out. The amount of the
primary αAlFeMnSi compound in the LM24 based RDC samples are summarised in Table 3, with
the volume percentage following a nearly linear correspondence to the Fe content. Scattered
primary βAlFeSi phase was reduced to a disc-like morphology. Some of the primary βAlFeSi plates
are indicated in Fig. 2b, showing their different cross-sections. Being the same as in the HPDC
samples, there were secondary intermetallic phases present in the form of patchy worm-like
morphologies that were often connected (e.g. Fig. 2b). These patchy regions of secondary
intermetallic phases corresponded to the inter-dendritic regions of the secondary α-Al solid solution
phase.
EDX quantification showed that the chemical compositions of primary intermetallic phases in the
LM24 based alloys were largely independent of the Fe contents of the alloys and processing
conditions (HPDC or RDC). Typical compositions of the αAlFeMnSi and βAlFeSi phases are listed
in Table 2, with the cubic αAlFeMnSi in the LM24 based containing some Cu and Cr. Intermetallic
phases in the patchy worm-like form were found to be largely copper bearing and EDX
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quantification from sites with marginally sufficient sizes (~ 1µm) showed compositions close to
θ(Al,Si)2Cu (Fig. 2b), with larger scatter in the quantification values due to the limitation of phase
sizes (Table 2). Si appeared to exist in the Al-sublattice of the θAl2Cu phase.
3.2 Microstructures of the LM25 based alloys
3.2.1 Microstructures of HPDC samples of LM25 based alloys
The microstructures of the HPDC samples were characterised by a mixture of dendritic α-Al solid
solution phase, interdendritic regions, being similar to the cases of LM24 based alloys. Primary Febearing intermetallic phases were present in alloys containing ≥0.74 wt.% Fe. The morphologies of
primary intermetallic phases vary with the Fe content. Alloys of lower Fe contents (0.74 wt.%) only
contained a small amount of spherical αAlFeMnSi compounds. On the other hand, for alloys
containing 0.98 and 1.13 wt.%Fe, the primary intermetallic phase exhibits morphologies of long
needles (e.g. Fig 3a). The EDX results on the compositions of primary intermetallic phases are
listed in Table 2, with the Fe-bearing compounds of the compact and needle-shaped morphologies
being identified as the αAlFeMnSi and βAlFeSi resectively. The compositions of the Fe-bearing
compounds in this group of alloys are very similar to these in the LM24 based alloys (Table 2). The
interdendritic regions of the HPDC samples are characterised with typical eutectic structures of the
Al-Si based casting materials, with quantities of secondary βAlFeSi needles associated with the
interdendritic regions increasing with the Fe content.
3.2.2 Microstructures of RDC samples of LM25 based alloy
For the primary α-Al solid solution phase, there are remarkable differences in the microstructures
of the RDC samples from their HPDC counterparts, similar to the case of the LM24 based alloys.
Instead of being dendritic as in their HPDC counterparts, the primary α-Al solid solution phase in
the RDC samples is largely spherical in morphology. The secondary α-Al phase exhibited dendritic
morphologies with associated interdendritic eutectic microstructures. As an example, Fig. 3b shows
the microstructure of the RDC sample for the LM25 alloy containing 1.13wt.%Fe.
Furthermore, RDC was found to be able to modify the morphologies of the primary intermetallic
phases as well. Even for the monoclinic βAlFeSi, RDC was able to reduce the long-needle
morphology into significantly shorter plates or discs (e.g. comparing the morphologies of the
primary βAlFeSi phase in Figs. 3a and 3b). Secondary β AlFeSi needles were also present in the
interdendritic eutectic regions, and they were morphologically similar to those in the HPDC
samples.
3.3 Mechanical Properties
Mechanical properties of the RDC and HPDC samples of the LM24 based alloys are compared in
Fig. 4. Elongation for the LM24 based alloys decreases with increasing Fe contents for both RDC
and HPDC samples. By comparison, the elongation values for the RDC samples are consistently
higher than that of their HPDC counterparts. The yield strength for both the RDC and HPDC
samples increases with Fe content up to a maxima at 1.65wt.%Fe. The ultimate strength of the RDC
samples of the LM24 based alloys is superior in alloys of higher Fe contents, with the disparity
between the RDC and HPDC samples increasing with Fe levels.
Mechanical properties of the RDC and HPDC samples of the LM25 based samples are compared
in Fig. 5. It is seen that the elongation of RDC samples of the LM25 based materials with different
Fe levels are all effectively higher than these of their HPDC counterparts. Overall, elongation
decreases with increasing Fe contents for both RDC and HPDC samples. However, the
enhancement of elongation for the RDC samples is so effective that even for the 1.13 wt.%Fe, the
RDC material still has elongation (> 4%) comparable to commercial die-cast alloys of much lower
Fe contents. The yield strength of RDC samples increases with increasing Fe content, approaching
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that of the HPDC samples with increasing Fe content. By comparison, the yield strength of the
HPDC sample only increases slightly with increasing Fe content. While the ultimate strength of the
HPDC samples decreases with increasing Fe content, the corresponding RDC samples nearly
maintain constant ultimate strength for an Fe content up to 1.13 wt.%.
In summary, for both the LM24 and LM25 based alloys of different Fe contents, the mechanical
properties of RDC samples were found to be improved compared to their HPDC counterparts,
mainly in elongation and ultimate tensile strength. The yield strength was found to increase with the
Fe content up to a certain level depending on the contents of Mn and other alloying elements.
4. Discussion
4.1 Modification of microstructures
4.1.1 The effects of melt shearing
The first stage of RDC was performed in the slurry maker, which was set at a temperature in the
(Liquid + solid/s) phase field, as required by semi-solid processing. Micorstructural studies showed
that melt shearing was effective to spheriodise the primary α-Al solid solution particles that were
formed in the slurry maker. It is well known now that intensive forced convection in the melt helps
to remove the dendritic morphologies of solid solution phases, leading to spherical/globular
morphologies as observed in the primary α-Al solid solution in this work [17,19]. Monte Carlo
simulations showed that spheriodisation of the primary solid solution phase in the slurry maker is
attributed to the reduction of the diffusion and thermal fields at the solid/liquid interface, and the
rotation of the solid particles during solidification [19]. The reduction of diffusion and thermal
fields in front of the solid/liquid interface helps to eliminate constitutional undercooling and hence
reduce the tendency for the development of dendritic morphologies during solidification.
The results of this work show that a similar effect of melt shearing on the morphologies of
intermetallic phases also exists. For the Fe-bearing intermetallic phases, melt shearing also showed
considerable effects in reducing the tendency for the formation of faceted features (Note that
faceting is common for intermetallic compounds), such that the star-like or cross-like morphologies
of the cubic αAlFeMnSi compound could be turned into spherical appearances (Figs. 1 and 2).
Even for the monoclinic βAlFeSi phase, melt shearing was seen to reduce their usual long-needle
morphology into discs or short bars (e.g. Fig. 3). The effects of melt shearing on the effects of the
cubic αAlFeMnSi could be interpreted in a similar way to the solid solution phase, as the star-like
or cross-like features can be considered as a special case for dendritic growth. On the other hand,
the effects of melt shearing on the morphological modification of the monoclinic βAlFeSi are
interesting, since the formation of needle-shaped morphology is largely owing to significantly
different solid-liquid (S-L) interfacial energy and associated disparity in growth rates along
different crystallographic directions. The modification of aspect ratios for the βAlFeSi phase seems
to suggest that melt shearing could reduce the differences in S-L interfacial energies to slow down
preferential growth in the direction of the longitudinal needle-axis. This could occur by the
reduction in the diffusion layer thickness at the S-L interface and the associated disturbance of local
equilibrium ahead of the growing solid front. This is expected to lower the compositional
difference at the S-L interface, leading to lowered S-L interfacial energy at the fast-growing front of
the solid phase. Such an effect in reducing the diffusion layer thickness would be biased, as the
thicker diffusion layer at the faster growing S-L front would experience bigger shearing effect.
It was observed that melt shearing in the slurry maker helped to increase the quantity and number
density of the cubic αAlFeMnSi compounds at the expense of the monoclinic βAlFeSi (e.g. Figs. 1
and 2). In the LM24 based alloys of higher Fe contents (≥1.65wt.%), the observed quantities of
αAlFeMnSi was even higher than the equilibrium amounts, with little βAlFeSi formed as primary
intermetallic compound. It is possible that intensive melt convection in the slurry maker exerted
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biased effects on the nucleation kinetics for the αAlFeMnSi and βAlFeSi phases, such that
nucleation of the former is promoted at the expense of the latter. In this way, one can make the best
use of Mn for replacing the monoclinic βAlFeSi with the cubic αAlFeMnSi phase. Indeed, by
comparison, the critical Mn/Fe ratios for the elimination of primary βAlFeSi are effectively lower
for the RDC samples than for their HPDC counterparts.
Overall, melt shearing was more effective for modifying the morphology of the quaternary
αAlFeMnSi phase (cubic) than for the ternary βAlFeSi (monoclinic) phase. The effects of melt
shearing on the amounts of βAlFeSi were found to be different in the two groups of alloys. In the
LM25 based alloys (of lower Mn content), melt shearing only seemed to affect the aspect ratios, but
it had little effect on the quantities of βAlFeSi. On the other hand, in the LM24 based alloys, which
contain a higher level of Mn, primary βAlFeSi needles were eliminated from the RDC
microstructures for the Fe contents up to 1.65 wt.%. Even in the LM24_2.04wt.%Fe alloy, the
amount of βAlFeSi was significantly reduced through melt shearing (Fig. 2). The micrsotructural
evolution in RDC samples was, therefore, controlled by both the physical melt shearing and the
alloy chemistry. Melt shearing was effective in reducing the amount of primary βAlFeSi in the
semi-solid state when there was a sufficient level of Mn.
4.1.2 The effects of alloying
In order to understand the effects of alloying on solidification microstructures, we carried out
CALPHAD modelling of the multi-component Al-Cu-Fe-Mg-Mn-Si-Zn system. The low levels of
impurities of other transition metal elements, such as Cr and Ti, are not considered to be significant
to the thermodynamic behaviour of the alloys of interest for this work. Calculations showed that
the level of Zn and Mg, which have considerable equilibrium solubility in the α-Al solid solution,
had little effects on equilibrium phase boundaries with intermetallic compounds. The most
important elements are therefore those that have low equilibrium solubility in α-Al, i.e. Fe, Mn, Cu
and Si. We used the COST507 thermodynamic database for constituent binary and ternary alloy
systems [15]. In order to describe the Al-Si-Fe-Mn system for alloys based on Al-Si, we followed
Balitchev et al [14] in treating the αAlFeMnSi as a stoichiometric phase, except that we chose a
stoichiometry of lower Mn content to be in line with experimental findings in the Al-Si based alloys
[9] (see Appendix). Such a description was found to predict pseudo-binary Al-Si-Fe-Mn diagrams
that are in excellent agreement with the experimental data of Zakharov et al. [13].
The effects of Si, Mn and Cu on the selection of Fe-bearing primary compounds can be seen
clearly by the pseudo-ternary liquid surface projections calculated in this work (Fig. 6a). It is seen
that at each level of constant Si content, the pseudo-ternary (Al-xSi)-Mn-Fe liquid surface
projections (x = 8, 10, and 12 wt.% respectively) are divided into regions for different primary
solidification phases: the L + ( Al ) , L + αAlMnSi , L + αAlFeMnSi , and L + βAlFeSi regions. One
can avoid the primary βAlFeSi solidification either by lowering the Fe content, or by increasing the
Mn/Fe ratio. For example, the dotted arrow-headed line in Fig. 6a defines a critical Mn/Fe ratio for
replacing the βAlFeSi primary phase with the cubic Fe-bearing primary phases, αAlFeMnSi or
αAlMnSi. This critical Mn/Fe ratio increases with increasing Si content. The effect of Cu is seen to
shift the L + β AlFeSi region to lower Fe content, without much influence on the critical Mn/Fe
ratio.
We superimposed the two groups of alloys investigated in this work on Fig 6a, to show their
relative locations on the pseudo-ternary liquid surface projections. The LM25 based alloys can be
closely referred to the (Al-10wt.%Si)-Mn-Fe, and the LM24 based alloys to the (Al-9wt.%Si3wt.%Cu)-Mn-Fe pseudo-ternary liquid surface projections. Fig. 6a shows that for the LM25 based
alloys (Mn ~ 0.16wt.%), primary βAlFeSi solidification cannot be avoided when Fe content
exceeds 0.8 wt.%. This is consistent with the experimental observation of the HPDC samples,
which showed the presence of large βAlFeSi needles only in alloys of higher Fe contents (> 0.74
wt.% of the four alloys). On the other hand, all five alloys of the LM24 based are located in the
8
L + αAlFeMnSi primary solidification region, though alloys of the highest Fe contents are close to
the L → αAlFeMnSi + βAlFeSi eutectic valley. Experimental investigation showed that long
βAlFeSi needles were only observed to coexist with αAlFeMnSi primary particles when
Fe≥1.34wt.%.
A 590°C pseudo-ternary (Al-10wt.%Si)-Mn-Fe isothermal section is shown in Fig. 6b, to
compare microstructures of RDC samples that were sheared in the slurry-maker. Such a shearing
stage was largely equivalent to isothermal annealing, which was enhanced kinetically due to
intensive forced convection in the liquid phase [21]. The LM25 based alloys are located in the
L + αAlFeMnSi + βAlFeSi for Fe content ≥0.74 wt.%. This is consistent with the microstructures of
RDC samples where βAlFeSi was found to coexist with the cubic αAlFeMnSi primary phase for Fe
≥ 0.74 wt.%. On the other hand, the LM24 based alloys with ≥1.34 wt.% Fe are located in
the L + αAlFeMnSi + βAlFeSi of Fig. 6b. Microstructural investigation showed the presence of
βAlFeSi in alloys with Fe content ≥1.65 wt.%. This is attributable to the presence of Cu and Cr in
the alloys, as these elements are well known to help for stabilising the αAlFeMnSi phase with
respect to the βAlFeSi in a similar effect to Mn [1] (Note the enrichment of Cu and Cr in the
αAlFeMnSi phase, as shown in Table 2). This is clearly shown in Fig. 6a, which exhibits the
enlargement of the L + αAlFeMnSi region and the narrowing of the L + β AlFeSi region due the
presence of Cu.
The effects of the Mn content on the formation of Fe-bearing intermetallic compounds are further
demonstrated in Fig. 7 for the Al-Fe-Si-Mn alloys with two different levels of Mn contents, 0.16
and 0.30 wt.% respectively. It is noticed from Fig. 7, that for the Mn and Fe levels of interest, α-Al
starts to solidify at a temperature about 594 °C, and Si starts to solidify at about 575°C (indicated).
This is in excellent agreement with the corresponding temperatures determined experimentally
(594±2 and 574±2 °C respectively [13]).
The solidification path for a LM25 based alloy (containing 0.74 wt.% Fe), is shown in Fig. 8, in
order to demonstrate the major features of the solidification process and the associated changes of
liquid compositions. These paths were calculated under equilibrium and Scheil [23] conditions.
Due to the extremely low solubility of the alloying elements Fe, Mn and Si in α-Al, there is little
difference between corresponding paths calculated under the two conditions. It is seen that the
solidification of α-Al leads to enrichment of all three elements (Si, Fe and Mn) in the remaining
liquid. Mn will be largely consumed by the solidification of the cubic αAlFeMnSi. The
solidification of the βAlFeSi phase leads to the enrichment of Mn and Si in the remaining liquid. An
adequate level of Mn is thus necessary, in order to turn the primary βAlFeSi (monoclinic) into the
cubic αAlFeMnSi phase. The final stage of solidification of the alloys typically based on LM25 is
the multi-eutectic transformation L → ( Al ) + Si + αAlFeMnSi + βAlFeSi to generate the eutectic
structure mainly of Al/Si with some intermetallic compounds, being in agreement with the
microstructures observed in this work. Similar solidification paths for the compositionally more
complex LM24 based alloys can be calculated using the thermodynamic database of this work,
which show the formation of θ(Al,Si)Cu2 and Mg2Si phases to consume Cu and Mg in the later
stage of solidification.
4.2 Microstructure-properties relationship
RDC was effective in improving mechanical properties for both the LM24 and LM25 based alloys,
particularly in improving the ductility of the alloys, so that the elongation of the RDC samples was
consistently higher than that of their counterpart HPDC samples. This is attributed to the
morphological modification in the primary Fe-bearing intermetallic compounds. RDC introduced
overall superiority in the elongation largely due to its effective role on modifying the morphologies
and sizes of the primary intermetallic compounds, so that long-needle shaped βAlFeSi was either
eliminated or modified into less harmful shapes.
9
The ultimate tensile strength of the RDC samples was effectively improved with respect to their
HPDC counterparts in alloys of higher Fe contents, while RDC was less effective on yield strength.
Overall, for RDC samples, the elongation decreased with increasing Fe contents, accompanied by
an increase in yield strength. For the LM25 alloys that contain about 0.16 wt.%Mn, the yield
strength of RDC samples approached that of the HPDC samples in alloys of higher Fe content. For
the LM24 based alloys, yield strength experienced maximum at 1.65 wt.% Fe, beyond which, it
began to drop. Referring to the sample microstructures, the enhanced yield strength corresponded
to increased amounts of Fe-bearing intermetallic compounds. The increase in yield strength
accompanies decreasing elongation, as the added reinforcement due to the Fe-bearing compounds is
at the cost of the alloy ductility.
Low Fe tolerance has been a major problem for recycling Al alloys, as Fe pickup and the
associated appearance of large needles of low-symmetry Fe-bearing compounds such as βAlFeSi,
causes remarkable loss of alloy ductility. The Fe tolerance of Al-alloys is therefore largely
determined by the loss in ductility. If we use the elongation values of the HPDC samples that
contain 0.7wt.% Fe as reference, we found that to maintain the same elongation value, RDC
increased the Fe tolerance from 0.7 to 1.36 wt.% (an increase of over 90%) for the LM24 based
alloys. For the LM25 based alloys, the increase of Fe tolerance from 0.7 up to 0.92 wt% due to
RDC corresponds to an improvement of about 30%. Such significant improvements of Fe tolerance
would help to remove a major barrier for recycling Al-alloys.
5. Conclusions
The Fe tolerance of Al-Si based alloys is closely correlated to the morphologies of the primary Febearing intermetallic compounds, and the detrimental effect of Fe can be eliminated by the removal
of large needle-shaped primary βAlFeSi compound. Chemically, this could be realised by limiting
the maximum level of Fe impurity, or by alloying with elements such as Mn to replace the
monoclinic βAlFeSi with a cubic αAlFeMnSi phase.
The critical Mn/Fe ratio for the elimination of the primary βAlFeSi varies with alloy composition.
Thermodynamic modelling showed that the ratio increases with increasing Si content. Cu on the
other hand, enlarges of the L + αAlFeMnSi region and promotes αAlFeMnSi solidification.
RDC is effective in modifying morphologies of the primary phases that solidify in the slurry
maker. It is effective in spheriodizing both the α-Al solid solution phase and the cubic αAlFeMnSi
compound. Even for the monoclinic βAlFeSi, RDC is effective in turning the usually long-needle
morphology into less detrimental short bars or discs.
Intensive melt shearing in the semisolid state promotes the solidification of the cubic αAlFeMnSi
phase at the expense of the monoclinic βAlFeSi. The Fe tolerance of Al-Si based cast alloys can be
effectively improved by combining Mn alloying and melt shearing. The application of RDC will be
an effective way in combating the detrimental effect of Fe on mechanical properties and is expected
to play a significantly beneficial role in recycling of Al-alloys.
Appendix
The thermodynamic description of the αAlFeMnSi phase, as referred to SGTE data of pure
elements [24], is:
BCC
α
G αAlFeMnSi − 0.72 0G AlFCC − 0.11 0G Fe
− 0.06 0GMn
− 0.11GSidiamond = - 21350 + 3.2 T (J / mole_atoms)
10
References
[1] L.F. Mondolfo. Aluminum alloys: structure and properties, London, Butterworths 1976.
[2] L. Wang, M. Makhlouf and D. Apelian: Int. Mater. Rev. 1995; 40: 221.
[3] G.B. Winkelman, Z.W. Chen, D.H. StJohn, M.Z. Jahedi. J. Mater. Sci. 2004; 39: 519.
[4] J.Z. Yi, Y.X. Gao, P.D. Lee, T.C. Lindley. Mater. Sci. Eng. 2004; A386: 396.
[5] Villars P, Calvert LD (ed.), Pearson’s Handbook of Crystallographic Data for Intermetallic
Phases, 2nd ed., ASM International, Materials Park, OH, 1991.
[6] Al and Al-alloys, ASM specialty handbook, ed. J.R. Davis, ASM International 1993.
[7] Smithells Metals Reference Book, 8th edition, ed. WF Gale and TC Totemeier, Elsevier 2004.
[8] S.G. Shabestari, Mater. Sci. Eng. 2004; 383A: 289.
[9] G. Davignon, A. Serneels, B. Verlinden, L. Delaey, Met. Mater. Trans. 1996; 27A: 3357.
[10] H.W.L. Phillips, P.C. Varley. JIM 1943; 69: 317.
[11] G. Phragmén, JIM 1950; 77: 489.
[12] J.G. Barlock, L.F. Mondolfo. Z. Metallk. 1975; 66: 605.
[13] A.M. Zakharov, I.T. Gul’din, A.A. Aenol’d, Yu A. Matsenko. Izv. Akad. Nauk SSSR Metall.
1989; 4: 214.
[14] E. Balitchev, T. Jantzen, I. Hurtado, D. Neuschütz. Computer Coupling of Phase Diagrams
and Thermochemistry 2003; 27: 275.
[15] I. Ansara, A.T. Dinsdale and M. H. Rand (eds.), COST 507-final report: Thermodynamic
Database for Light MetAl-alloys, Vol.2, European Communities, Brussels 1998.
[16] Z. Fan, X. Fang and S. Ji. Materials Science and Engineering A 2005; 412: 298.
[17] Z. Fan, Int Mater Rev 2002; 47: 49.
[18] A. Das and Z. Fan: Material Science and Technology 2003, 19: 573.
[19] A. Das, S. Ji and Z. Fan. Acta Mater. 2002; 50: 4571.
[20] Z. Fan, G.J. Liu. Acta Mater. 2005; 53: 4345.
[21] Y.Q. Liu, A. Das, Z. Fan. Mater. Sci. Tech. 2004; 20: 35.
[22] ASTM standard B557: ‘Standard methods of tension testing wrought and cast aluminum and
magnesium alloy products’, Annual book of ASTM standard, vol. 02.02, 1993.
[23] E. Scheil, Z Metallkde 1942; 34: 70.
[24] A.T. Dinsdale, CALPHAD 1991; 15(4): 317.
11
Table 1 Alloy compositions in wt. %.
Alloys
Si
Fe
Cu
Mn
Mg
Zn
Cr
Ti
Al
wt. %
LM24
9.30
0.70
0.97
1.34
1.65
2.04
3.25
0.30
0.26
1.58
0.04
0.05
balance
LM25
10.30
0.46
0.74
0.98
1.13
0.30
0.16
0.30
0.25
--
0.05
balance
Table 2 Compositions of intermetallic phases in at. %, from quantitative EDX analysis, with
compositions below statistical errors ignored.
Phases
Si
Fe
Cu
Mn
Mg
Zn
Cr
Ti
Al
at. %
αAlFeMnSi
θAl2Cu
βAlFeSi
9.7±0.2
5.12±2.5
18.7±0.5
12.1±0.9
-13.9±0.5
0.7±0.2
16.4±5.2
--
LM24 based:
3.2±0.5
---0.6±0.1
--
----
0.6±0.3
---
----
balance
balance
balance
βAlFeSi
αAlFeMnSi
18.4±0.3
10.8±0.5
14.2±1.5
12.5±0.9
---
LM25 based:
0.6±0.1
-3.3±0.5
--
---
---
---
balance
balance
Table 3 Amount of primary αAlFeMnSi particles in the LM24 based RDC samples.
Fe content (wt. %)
Volume percentage
Particle density (mm-2 )
0.97
1.34
1.65
2.04
0.5
1.3
3.0
4.3
5
25
67
77
12
αAlFeMnSi
β AlFeSi
(a)
α-Al
E
β AlFeSi
(b)
Fig. 1 Optical micrographs for the HPDC LM24_2.04 wt% Fe alloy (pouring temperature 640 °C).
(a) General view to show the presence of primary intermetallic phases – the compact αAlFeMnSi
quaternary compound and the needle-shaped βAlFeSi ternary compound. (b) Higher magnification
to show the dendritic α-Al (light), the interdendritic eutectic regions (indicated as “E”) and the large
primary βAlFeSi needles.
13
(a)
αAlFeMnSi
α-Al
50 µm
βAlFeSi
αAlFeMnSi
βAlFeSi
(b)
θAl2Cu
50 µm
Fig. 2 RDC LM24_2.04wt%Fe alloy: (a) Optical micrograph for the overall distribution of
primary α-Al and αAlFeMnSi. The inset is an enlarged view of the grey background, showing the
dendritic secondary α-Al and interdendritic eutectic structures. (b) Backscattered electron image
(BSI) showing the distribution of intermetallic compounds. Different crosss-sections for primary
βAlFeSi compound are indicated in (b). The patchy intermetallic phase in (b) is mainly the Cu
bearing aluminide θAl2Cu, with Si substituting some Al sublattices. The α-Al and Si phases cannot
be distinguished from each other by BSI due to their similar atomic mass.
14
(a)
βAlFeSi
α-Al
(b)
α-Al
βAlFeSi
αAlFeMnSi
20 µm
Fig. 3 Optical micrographs of the LM25_1.13wt%Fe alloy: (a) HPDC sample, showing the
presence of needle-like primary βAlFeSi compounds. (b) RDC sample, with the inset of same
magnification showing different sections of the βAlFeSi phase.
15
UTS (MPa)
340
320
300
280
260
240
YS (MPa)
180
160
140
120
Elongation (%)
5.0
4.0
3.0
RDC
HPDC
2.0
1.0
0.0
0.5
1.0
1.5
2.0
2.5
Fe content (wt.%)
Fig. 4 Comparison of Fe effect on mechanical properties of RDC and HPDC samples of the LM24
based alloys.
16
UTS (MPa)
300
290
280
270
260
250
YS (MPa)
160
140
120
100
Elongation (%)
12
10
8
6
4
RDC
HPDC
2
0
0.2
0.4
0.6
0.8
1.0
1.2
Fe content (wt.%)
Fig. 5 Comparison of the effect of Fe on mechanical properties of RDC and HPDC samples of the
LM25 based alloys.
17
2.0
1.8 BETA
1.4
Fe ( wt% )
ALFEMNSI_A
L + βAlFeSi
1.6
LM24 based
LM25 based
--- 8Si
--- 10Si
--12Si
--- 9Si-3Cu
L + αAlFeMnSi
1.2
1.0
0.8
0.6
L + αAlMnSi
0.4
AL
0.2
0
ALMNSI_ALPHA
L + ( Al )
0
0.2
0.4
0.6
0.8
1.0
1.2
1.4
1.6
1.8
2.0
Mn ( wt% )
(a)
2.0
(Al
l)
(A
lFe
Mn
Si
+
+α
L+
αA
lFe
Mn
L+
Si
αA
1.2
L+
β
L
+
β
1.4
(V)
Al F
eM
nS
1.6
Fe ( wt% )
590 C (10Si)
(IV)
)
(III)
i+
1.8 (II)
1.0
0.8
0.6
L+ αAlFeMnSi + αAlMnSi + (Al)
0.4
0.2
0
L+ αAlMnSi
L+ ( Al )
(I)
0
0.2
(VI)
0.4
0.6
0.8
1.0
1.2
1.4
1.6
1.8
2.0
Mn ( wt% )
(b)
Fig. 6 Phase diagrams calculated in this work. (a) Liquid surface projections for the quaternary AlxSi-Fe-Mn (x = 8, 10 and 12 wt% respectively), and for Al-9wt%Si-3.5wt%Cu-Fe-Mn, showing the
effects of Si and Cu contents on primary solidification phases. (b) 590°C isothermal section of
quaternary Al-10wt%Si-Fe-Mn alloys. The borders with the β phase are highlighted in (b). The
superimposed triangles and squares represent the Fe and Mn levels of the LM25 and LM24 based
alloys respectively.
18
700
(a)
680
10Si-0.16Mn
660
640
L + βAlFeSi
L + ( Al ) + αAlFeMnSi
T (C)
620
L + αAlFeMnSi + βAlFeSi
600
(Al)
L + ( Al )
580
L + ( Al ) + αAlFeMnSi + βAlFeSi
Si
560
( Al ) + Si + αAlFeMnSi + αAMnSi
540
( Al ) + Si + αAlFeMnSi + βAlFeSi
520
500
2.0
1.5
1.0
0.5
0
( Al ) + Si + αAlMnSi
2.5
3.0
Fe ( wt% )
700
(b)
680
10Si-0.3Mn
L + βAlFeSi
660
640
L + αAlFeMnSi
L + ( Al ) + αAlFeMnSi
T (C)
620
L + αAlFeMnSi + βAlFeSi
600
(Al)
L + ( Al )
580
L + ( Al ) + αAlFeMnSi + βAlFeSi
Si
560
( Al ) + Si + αAlFeMnSi + αAMnSi
( Al ) + Si
+ αAMnSi
540
( Al ) + Si + αAlFeMnSi + βAlFeSi
520
500
0
0.5
1.0
1.5
2.0
2.5
3.0
Fe ( wt% )
Fig. 7 Pseudo-binary phase diagrams of the Al-Fe-Mn-10wt%Si alloys with Mn being fixed at (a)
0.16wt%, and (b) 0.3 wt%. The starting temperatures for the solidification of the α-Al and Si phase
are indicated respectively. Mn is seen to enhance the solidification of primary αAlFeMnSi, by
extending the (L+αAlFeMnSi) to higher Fe level.
19
o
T ( C)
Sch
eq
595
590
585
580
575
Weight fraction in liquid
0.00
0.125
0.120
0.115
0.110
0.105
0.100
0.05
0.10
0.15
L --> (Al)
fs
L --> (Al) + α AlMnFeSi
0.25
Si
L --> (Al) + β AlFeSi
L --> (Al) + α AlMnSi + Si
0.00
0.0018
0.0017
0.20
0.05
I
0.10
0.15
0.20
II III
0.25
IV
Mn
0.0016
0.00
0.0080
0.0075
0.0070
0.0065
0.0060
0.0055
0.05
0.00
0.05
0.10
0.15
0.20
0.25
0.15
0.20
0.25
Fe
0.10
Mole fraction of solid
Fig. 8 Solidification paths of the LM25_0.74wt%Fe alloy, showing little difference between the
paths calculated under the equilibrium and Scheil solidification conditions. It is seen that during
solidification, Mn is mainly consumed by the αAlFeMnSi phase, and Fe is mainly consumed by
βAlFeSi. The consumption of Fe by the αAlFeMnSi phase is counterbalanced by the rejection of
Fe from the α-Al phase.
20