Aluminium phosphide as a eutectic grain nucleus in J Electron Microsc (Tokyo)-2004-Nogita-361-9.pdf

7/28/2019 Aluminium phosphide as a eutectic grain nucleus in J Electron Microsc (Tokyo)-2004-Nogita-361-9.pdf

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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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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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(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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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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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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[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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growth of Al dendrite. (b) AlP particles act as nucleation sites for pol-

yhedral silicon crystals from which eutectic grains grow.


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Aluminium phosphide as a eutectic grain nucleus in J Electron Microsc (Tokyo)-2004-Nogita-361-9.pdf