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Journal of Electron Microscopy53(4): 361369 (2004) Japanese Society of Microscopy
doi: 10.1093/jmicro/dfh048
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Full-length paper
Aluminium phosphide as a eutectic grain nucleus inhypoeutectic Al-Si alloys
Kazuhiro Nogita1,*, Stuart D. McDonald1, Katsuhiro Tsujimoto2, Kazuhiro Yasuda3 and Arne
K. Dahle1
1Division of Materials Engineering, University of Queensland, St Lucia 4072, Australia, 2Research Laboratory forHigh Voltage Electron Microscopy, Kyushu University, Japan and 3Department of Applied Quantum Physics andNuclear Engineering, Kyushu University, Japan*To whom correspondence should be addressed. E-mail: [email protected]
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Abstract Aluminium phosphide (AlP) particles are often suggested to be the nuclea-tion site for eutectic silicon in Al-Si alloys, since both the crystal structureand lattice parameter of AlP (crystal structure: cubic F43m; lattice parame-
ter: 5.421 ) are close to that of silicon (cubic Fd3m, 5.431 ), and themelting point is higher than the Al-Si eutectic temperature. However, thecrystallographic relationships between AlP particles and the surroundingeutectic silicon are seldom reported due to the difficulty in analysing theAlP particles, which react with water during sample preparation for polish-ing. In this study, the orientation relationships between AlP and Si are ana-lysed by transmission electron microscopy using focused ion-beam millingfor sample preparation to investigate the nucleation mechanism of eutecticsilicon on AlP. The results show a clear and direct lattice relationshipbetween centrally located AlP particles and the surrounding silicon in thehypoeutectic Al-Si alloy.
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Keywords Al-Si alloy, eutectic, solidification, nucleation, aluminium phosphide,focused ion beam
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Received 2 March 2004, accepted 21 May 2004
Introduction
Aluminium-silicon (Al-Si) foundry alloys are popular because
of their good castability, surface finish and resistance to corro-
sion, coupled with their high strength-to-weight ratio. Typi-
cally, ~50100% of the volume of an Al-Si casting alloy is a
eutectic mixture of aluminium and silicon. Solidification of
the eutectic is one of the last reactions to occur during cooling
and this is where most casting defects, such as porosity and
hot tears, are formed. These defects reduce the mechanical
properties, compromise the ability of the casting to contain
liquid or gas under pressure, interfere with coating and
machining operations and can detract from the appearance of
the product, in addition to causing a high reject rate. Although
it is frequently acknowledged that the solidification of the
eutectic is critical for porosity formation [14], there is cur-
rently a lack of understanding about how the eutectic evolves
(particularly nucleation) and how it can be controlled.
It has only recently been determined, using the technique of
electron backscattered diffraction (EBSD), that there are three
eutectic solidification modes (or macroscopic growth pat-
terns) that operate in Al-Si alloys [4]. Figure 1 shows an illus-
tration of the three modes, with supporting evidence of their
existence from the optical microscopy of samples quenched
during eutectic solidification. These are (a) nucleation on or
adjacent to primary aluminium dendrites, (b) independent
heterogeneous nucleation of eutectic grains in interdendritic
spaces and (c) growth of the eutectic solidification front oppo-
site to the thermal gradient. More than one growth mode may
operate at any given time, although the addition of certain
elements can cause a preference for one of the above mecha-
nisms. It is likely that the different eutectic solidification
modes arise due to the activation and/or inactivation of vari-
ous nuclei. Ideally, it is desirable to control both the silicon
morphology and the eutectic growth mode independently.
This will allow both the macro- and microstructures of the
eutectic to be tailored as required.
Figure 2 is a micrograph of an unmodified Al-Si alloy that
has been quenched during eutectic solidification to identify
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J O U R N A L O F E L E C T R O N M I C R O S C O P Y, Vol. 53, No. 4, 2004362
(a) (b) (c)
Mould wallPrimary dendrite
Al-Si eutectic
100m 200m 200m
Fig. 1 Eutectic growth modes. (a) Nucleation on or adjacent to primary aluminium dendrites, (b) independent heterogeneous nucleation of
eutectic grains in interdendritic spaces and (c) growth of the eutectic solidification front opposite to the thermal gradient.
(a) (b)
Fig. 2 Optical micrographs of quenched sample during eutectic solidification at (a) low magnification and (b) higher magnification with nuclei
particle centre of eutectic silicon.
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K. Nogitaet al. Aluminium phosphide as a eutectic grain 363
eutectic grains. Many of the eutectic grains contain centrally
located particles within the silicon phase. It is tempting to
conclude that such particles may have been a nucleus for the
eutectic grain. In Al-Si alloys, it has often been suggested that
particles similar in appearance to this one are AlP, because AlP
and Si have near identical lattice parameters, theoretically
making it an ideal nucleus. Conventional energy-dispersive
spectroscopy (EDS) analysis (Figure 3) does show the pres-
ence of Al and P; however, oxygen (O) is almost always
present in large quantities [5], raising the possibility that the
particle is actually aluminium phosphate (AlPO4), which is a
common contaminant of aluminium-based melts. Unfortu-
nately, conclusive proof of nucleation cannot be obtained byobserving an intimate physical relationship between phases
such as that seen in Figure 3, but is reliant on crystallographic
evidence provided by transmission electron microscopy
(TEM). This makes it difficult to identify a nucleus with confi-
dence for two reasons: firstly, a great number of TEM samples
would need to be prepared before a suitable particle could be
found within the prepared region and, secondly, the sample
can be damaged during conventional preparation procedures.
The technique of focused ion beam (FIB) milling and in situ
sample manipulation allow eutectic grains to be excavated
and removed for TEM analysis of potential nuclei. This study
uses this advanced analysis technique to locate and identify
the composition of particles that are potential eutectic grain
nuclei and test for the existence of an orientation relationship
between these particles and surrounding phases.
MethodsAluminium-silicon alloys of nominal composition (Al
10mass% Si) were used for experimentation. The composition
of the base alloy is given in Table 1. The alloys were produced
by placing ~1 kg commercial purity aluminium and silicon in
a clay-graphite crucible and heating in an electric resistance
O
P
A
Si
Al
Distance (m)
Counts(arbitraryunits)
Counts(arbitraryunits)
0 10 20 30 40
Fig. 3 The EDS line scan results showing the presence of Al, P and O in the centrally located nucleant particle.
.....................................................................................................................................................................................................................
Table 1. Chemical composition of the sample (wt.%)
....................
Al.............
Si..............
Cu..............
Fe.................
Mg.................
Zn...................
Cr.................
Ni.................
Mn.................
Ti...................
Sr...................
Zr...............
P
Balance 9.77 0.83 0.11
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J O U R N A L O F E L E C T R O N M I C R O S C O P Y, Vol. 53, No. 4, 2004364
furnace to a temperature of 760C. Phosphorus additions were
made in the form of a Cu-P master alloy to raise the final
phosphorus content to an analysed level of 30 p.p.m.
Metallographicsamplesweretakenbypouringthemeltintoa copper mould at room temperature.
Samples were prepared for microscopy using standard
techniques, with the final polishing stage being produced by
0.05 m colloidal silica. Samples were etched for 60 s at room
temperature in a mixture of 60 ml water, 10 g sodium hydro-
xide and 5 g potassium ferricyanide (Modified Murakami
reagent). Optical micrographs were taken of both unetched
and etched samples.
Eutectic grains were examined using scanning electron
microscopy and those grains that contained visible potential
nuclei (i.e. centrally located particles) were selected for further
preparation using FIB milling. It is only recently that the
advanced technique of FIB milling combined with in situ sam-
ple micromanipulation has become available to selectively
extract a TEM sample from any desired region within a larger
sample. The FIB sample preparation was performed with a
Hitachi FB-2000A with 30 kV Ga liquid metal ion source. The
procedure that was used is outlined in Fig. 4ag, which shows
the sample prior to milling (a), after being protected with a W
coating (b), after excavation using a 30 kV Ga source (c), being
tilted to 45 to allow the sample to be completely cut (d), dur-
ing manipulation with a microneedle after FIB milling (e),
after being attached to a TEM sample grid and further thinned
(30 kV Ga) (f), and ready for TEM analysis (g). The TEM
observationswere conducted using a 200 kV Technai20TEM and included the acquisition of micro-selected area
diffractions (SAD) taken from an area of ~60 nm in diameter,
small probe convergent-beam electron diffraction patterns
(a) (b)
(c)(d)
Al
Si
Nucleant
Fig. 4 (ad)
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K. Nogitaet al. Aluminium phosphide as a eutectic grain 365
(CBED) taken from an area of ~10 nm in diameter and
energy-dispersive X-ray analysis.
Results and discussionThe crystal structure and lattice parameters of the close-
packed planes of Si, AlP and AlPO4
are shown in Table 2. From
this table it is apparent that there is minimal mismatch
between the AlP and the Si phases (
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J O U R N A L O F E L E C T R O N M I C R O S C O P Y, Vol. 53, No. 4, 2004366
the nuclei must also be present as a solid (although not neces-sarily thermodynamically stable) at the appropriate tempera-
ture and composition of the liquid, and also with a suitable
size distribution. The nuclei must also be capable of being wet-
ted by the liquid and should not be consumed or enveloped by
reactions that have occurred earlier in solidification. For these
reasons it is not without pitfalls to assume AlP is the nucleus
for eutectic solidification solely because of minimal lattice
mismatch between AlP and Si. This is compounded by the fact
that oxygen has been detected along with phosphorus in large
quantities in suspected nuclei and that AlPO4
is a common
refractory binder used in aluminium production. Experimen-
tal determination is therefore the best way to determine the
composition of many nuclei.
Figure 5 is a micrograph of a cross-sectioned sample pro-
duced using the FIB. The sample contains eutectic Al and Si
and a third phase that has possibly acted as a nucleus for the
surrounding silicon. From the micrograph it appears the
nucleus is not a discrete, compact particle, but is probably a
convoluted morphology (connected in three dimensions) or
alternatively one of several particles. Figure 6 shows a slightly
higher magnification imagealong with a micro-SAD pattern(area of ~60 nm in diameter) from the locations labelled A
(unknown particle), and small probe CBED patterns (area of~10 nm) from locations B (unknown particle) and C
(eutectic silicon). Location A lies within a particle that was
exposed to preparation using conventional polishing tech-
niques and the SAD pattern along with EDS results show that
this particle has an amorphous structure and contains Al, P
and O (Fig. 6a). In contrast, the CBED patterns from locations
B and C, which were deep to the polished surface and
exposed only by FIB milling, indicate crystalline structures. By
comparison of CBED patterns (obtained using an identical
angle of tilt after aligning the [011] pole in location C) for the
particle (location B) and the silicon crystal (location C)
there is clearly a common diamond structure. The EDS analy-
sis confirmed that the composition of the particle at location
B was AlP and no oxygen was present. The crystallographic
orientation relationship between the AlP (Fig. 6b) and neigh-
bouring eutectic Si (Fig. 6c) is identical, conclusively proving
that the AlP particle acted as an epitaxial nucleation site for
the eutectic silicon.
The above results confirm that AlP is the nucleation site for
eutectic silicon. Furthermore, it is tempting to conclude that
reactions occurring during conventional sample preparation
methods disrupt the crystalline nature of any exposed regions
Si
Al
Nucleant particle
Fig. 5 The FIB sample of eutectic Si in hypoeutectic Al-Si alloy.
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K. Nogitaet al. Aluminium phosphide as a eutectic grain 367
of AlP. The oxygen that is present in exposed particles (see, for
example, Fig. 3) is likely to be an artefact of the polishing
processes and does not indicate the presence of AlPO4.
Although FIB milling allowed the selective preparation of
TEM samples and avoided damage of the sample compared
with conventional procedures, it was not without disadvan-
tages. The first problem was due to a contamination of the sur-
face with W during the sample excavation and milling stages.
The second problem arose because the minimum sample
thickness achievable is limited to ~200 nm. This limited the
quality of the high-resolution lattice images that were obtain-
able from the sample and prevented the direct measurement
of lattice parameters.
The identification of AlP as a nucleation site for eutectic
silicon supports a model of eutectic nucleation that is strongly
influenced by impurities in the melt. A probable nucleation
sequence is described with the aid of Figure 7. During the first
stage of solidification in hypoeutectic Al-Si alloys, dendritic
growth occurs and phosphorus and silicon will segregate from
the dendriteliquid interface due to their low solubility in
solid aluminium. If sufficient phosphorus is present, AlP will
form at the dendriteliquid interface due to the increased sol-
ute content, as shown schematically in Fig. 7a. Alternatively,
the AlP particles may be stable in the melt and simply be
pushed ahead of the dendriteliquid interface during solidifi-
cation. Each of these particles can act as a nucleus for a silicon
crystal and the commonly observed polyhedral shape of the
silicon crystal, resembling that of primary silicon crystals, may
be adopted because of the locally high silicon concentration at
the interface. Eutectic solidification commences from each
growing polyhedral crystal (Fig. 7b).
Previous results have shown that the eutectic aluminium in
commercial unmodified alloys has an identical crystallo-
graphic relationship to the surrounding aluminium dendrites
A B C
Fig. 6 TEM image and corresponding (A) electron diffraction patterns from Al-P-O precipitate on polished surface showing amorphous struc-
ture, (B) small probe CBED pattern from AlP showing [011] pole, and (C) small probe CBED pattern from Si showing [011] pole. CBED patterns
for (B) and (C) were obtained using an identical tilt angle.
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J O U R N A L O F E L E C T R O N M I C R O S C O P Y, Vol. 53, No. 4, 2004368
[4,6]. While this may seem in conflict with nucleation of
eutectic silicon on AlP particles within the melt, this is not
necessarily so. Instead, it appears that the nucleation and
growth of the eutectic aluminium and silicon phases occur
somewhat independent from one another. When eutectic
silicon nucleates on AlP particles near the dendrite liquid
interface, the adjacent melt becomes enriched in aluminium
because of the almost zero solubility of aluminium in silicon.
Renucleation of aluminium is not necessary to allow coupled
eutectic growth, due to the proximity of the existing alumin-
ium dendrites. Instead, the eutectic aluminium is likely to
simply grow from the dendrites to surround the silicon phase.
A schematic representation of the complete nucleation
sequence for the Al-Si eutectic is shown in Fig. 7.
Cooling curve traces of Al-Si alloys further support the
above observation of the very efficient nucleation of silicon on
AlP particles [5]. These cooling curves show virtually no
undercooling and recalescence with the presence of phospho-rous in the melt, and the effect of phosphorous is evident even
at quite low concentrations (e.g. in the order of 10 p.p.m.).
Observations of quenched samples have also shown that the
nucleation frequency is relatively high in alloys containing
added phosphorous, and that addition of elements such as Sr
and Na (which are introduced to modify the silicon from
coarse plates to a fine fibrous morphology) causes a dramatic
decrease in nucleation frequency [4]. The exact mechanism by
which these elements cause a decrease of nucleation fre-
quency is unknown, but it has been speculated it may be a
result of deactivation of the AlP nucleants through coverage
by intermetallic phases and/or alterations of surface energy.
The high nucleation frequency observed with AlP is further
stimulated by the strong segregation of the eutectic alumin-
ium from silicon. This constitutional effect will strongly sup-
port the stability, survival and further growth of the silicon.
The success of combining the techniques in this work to
study potential nucleant particles could stimulate similar
studies of nucleation-related phenomena for a range of sys-
tems. For example, it would be most interesting to use a simi-
lar approach to investigate the nucleation of primary
aluminium crystals, in combination with other experiments.
Such studies could investigate the structure, properties and
composition of nucleant particles, perhaps adding more sub-
stantial results to resolve discussions regarding the role of
borides (AlB2, TiB
2and mixed borides) and aluminides (Al
3Ti)
in the grain refinement of aluminium alloys. Furthermore,
further evidence may be obtained to substantiate the recent
model by Greer et al. [7] about the role of nucleant particle size
in nucleation. Their model proposes that it is the larger rangeparticle sizes in the nucleant particle size distribution that
become activated on solidification and that it is the undercool-
ing (through the cooling rate and recalescence) that controls
the fraction of the particle size distribution that becomes acti-
vated as nucleation sites in a given alloy with a certain nucle-
ant particle distribution. However, the approach taken in this
paper is certainly not limited to the study of liquidsolid trans-
formations. For example, this approach could prove itself as a
powerful method to study lattice relationships in three dimen-
sions for any phase transformation and nucleant. However, it
is also important to realize that it is likely that the method
requires further evidence, such as from combining it with
other techniques, e.g. thermal analysis methods, quenching
and statistical analysis, to provide for complete analyses of the
phenomena governing the transformations at hand.
Concluding remarks
The technique of FIB milling along with micromanipulation
was used to create TEM samples from selected areas within
the microstructure of hypoeutectic Al-Si alloys. Selective sam-
ple preparation allowed TEM analysis of suspected nuclei that
would not have been possible with conventional sample prep-
aration. Using TEM analysis techniques, including micro-SAD
and CBED, it is shown experimentally that AlP particles andsurrounding eutectic silicon share a common orientation and
it is confirmed that AlP nucleates eutectic silicon. A likely
solidification sequence for eutectic solidification in unmodified
hypoeutectic Al-Si alloys is proposed. Conventional sample
preparation is shown to introduce a significant amount of
oxygen and damage the natural crystallography of the AlP
phase. The analysis technique detailed is ideally suited to
identifying potential nuclei in many alloys where a low
number of particles or high reactivity of the sample may limit
the use of conventional TEM samples.
AcknowledgementsThis research is supported by the Kyushu University High Voltage Electron
Microscopy Program under the Nanotechnology Support Project of the
Ministry of Education, Culture, Sports, Science and Technology (MEXT),
Japan (project ID: T-Kyudai-H15-013). The authors also acknowledge
financial support from the Center for Microscopy and Microanalysis
(CMM) and the University of Queensland.
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