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Materials Transactions, Vol. 52, No. 6 (2011) pp. 1308 to 1315
#2011 The Japan Institute of Metals
EXPRESS REGULAR ARTICLE
Prediction of Solidification Paths in Al-Si-Fe Ternary System
and Experimental Verification: Part II. Fe-Containing Eutectic Al-Si Alloys
Sanghwan Lee, Bonghwan Kim and Sangmok Lee
Liquid Processing & Casting Technology R&D Group, Korea Institute of Industrial Technology,
7-47 Songdo-Dong, Yeonsu-Gu, Incheon 406-840, Korea
The effects of Fe content and cooling rate on the solidification path and formation behavior of the Al5 FeSi () phase in Fe-containing
eutectic Al-Si alloys were studied based on thermodynamic analysis and pertinent experiments. To predict solidification path in Fe-containing
eutectic Al-Si alloys, the thermodynamic calculations in the Al-Si-Fe ternary system, including the high Fe region, were systematically
performed using the Thermo-Calc program and updated database. To experimentally verify the predicted results, designed two eutectic Al-Si
alloys with different Fe levels were solidified under slowly- and rapidly-cooled conditions, respectively. The cooling curves of the solidified
alloys were recorded by thermal analysis. Microstructures of the casting samples were studied by the combined analyses of optical microscopy
(OM) and scanning electron microscopy (SEM). For the slowly-cooled condition with the high Fe level, the primary phase was mainly formed
by the quasi-peritectic reaction of L þ ! þ (656 C). For the rapidly-cooled condition with high Fe level, the primary phase was mainly
formed by the reaction of L ! (583649 C). [doi:10.2320/matertrans.M2010423]
(Received December 13, 2010; Accepted March 10, 2011; Published May 18, 2011)
Keywords: eutectic aluminum-silicon alloys, casting, solidification path, intermetallic compound, iron content, cooling rate
1.
Introduction
Eutectic Al-Si alloys are a wide range of useful materials
in different fields of industry where high strength to weight
ratio, castability, and wear resistance is required.1–5) Si,
which is very chip as a raw material, is good and desirable
element in Al casting alloys because it decreases melting
temperature, increases melt fluidity and reduces thermal
expansion coefficient (related to the solidification contraction).3,4) For these reasons, eutectic Al-Si alloys have
excellent casting properties. Eutectic Al-Si alloys also have
a good abrasion resistance in accordance with forming a large
amount of hard (Si) phase due to the negligible solubility of
Si in Al alloys.4) Moreover, the overall weight of the cast
component adopting eutectic Al-Si alloys is very light
because the density of the alloys decreases almost linearly
with increasing Si content due to low density (2.33 g/cm3 )
of Si.5)
Due to the high liquid solubility and the low solid
solubility of Fe in Al-Si alloys, Fe has various ways to come
into the molten these alloys and promotes to form various
Fe-rich intermetallic phases such as Al8 Fe2 Si () and
Al5 FeSi ().6–8) For these reasons, Fe is always present in
commercial Al-Si alloys and has been regarded as the most
harmful element degrading mechanical properties of these
alloys.8–14)
Although many studies for various and complex phase
formation behaviors in the Al-Si-Fe system15–24) have
continued to be performed during the past decades, there
are a lack of precise information and comprehensive understanding in high Fe region of the ternary system. The main
purposes of the present study are to predict the solidification
path and phase formation based on the thermodynamic
analysis and to verify the predicted results experimentally.
The present paper reports the effects of Fe content and
cooling rate on those phase formation sequence in Fecontaining eutectic Al-Si alloys. The paper also reports the
Table 1 Nominal and actual chemical compositions of the casting alloys.
Chemical composition (mass%)
Specimen
Alloy 1
Alloy 2
Si
Fe
Other
Al
Nominal
12.40
0.30
—
Bal.
Actual
12.40
0.34
0.05
Bal.
Nominal
12.40
8.00
—
Bal.
Actual
12.35
8.17
0.39
Bal.
effect of Si content through comparing the results investigated in Part I.
2.
Materials and Methods
Two eutectic Al-Si alloys, with much the same Si level
(12.4 mass%) but different Fe levels, were prepared from
high purity Al (99.8 mass%), Al-15 mass%Si and Al75 mass%Fe master alloys. The experimental methods
including the two cooling conditions and the analysis
methods such as thermodynamic, microstructural and thermal analysis are essentially the same as in Part I.
Table 1 shows the analyzed compositions of the two
casting alloys by spark emission spectroscopy (SES, OBLF
QSN-750). For the precise thermodynamic analysis in high
Fe region, a thermodynamic database for Thermo-Calc was
correctly updated and revised using the collected up-to-date
references.23–27) For the investigation of slow solidification
by furnace cooling, experimental verification was performed
for the solidification path predicted by the thermodynamic
analysis. For the investigation of rapid solidification by ice
water quenching, the effect of cooling rate on the formation
behavior of phase was examined. The various intermetallic
compounds referred to in this study were identified by the
combined use of back-scattered electron (BSE) imaging
mode and energy dispersive X-ray spectroscopy (EDS,
EDAX Apollo SDD series) of scanning electron microscopy
(SEM, FEI Quanta-200F).
Prediction of Solidification Paths in Fe-Containing Eutectic Al-Si Alloys
(a)
14
0
95
(b)
°C
14
95
12
°C
12
ζ
ζ
0°
90
°C
00
9
C
10
10
C
0°
8
85
6
80 0
P1
°C
°
750
4
γ
θ
6
°C
800
α
U1 P2
650 °C
U2
(Al)
U3
600 °C
β
FeCR
2
4
6
δ
E1
8
10
12
θ
P1
°C
750
α
700 °C
2
U1 P2
650 °C
U2
U4
(Si)
0
0
85
4
γ
C
0°
8
C
700 °C
2
Fe (mass%)
Fe (mass%)
0
1309
14
Si (mass%)
U3
(Al)
600 °C
2
4
U4
E1
0
0
δ
β
6
8
10
12
(Si)
14
Si (mass%)
Fig. 1 Calculated liquidus projection and isotherms in the Al-rich corner of the Al-Si-Fe system and the predicted equilibrium
solidification paths; (a) alloy 1 and (b) alloy 2.
Fig. 2 Reaction sequences predicted under the assumption of equilibrium solidification and corresponding changes in liquid composition;
(a) alloy 1 and (b) alloy 2.
3.
Results
3.1 Solidification path
The solidification paths of two eutectic Al-Si alloys with
different Fe levels were predicted by assuming that solidification occurs under full equilibrium condition. The
predicted solidification paths of the two alloys are superimposed on the calculated liquidus projection of Fig. 1. As a
summary of solidification behaviors, the predicted reaction
sequences of two alloys and the corresponding calculated
changes in liquid composition are indicated in Fig. 2. For
alloy 1, the first solid phase (Al) begins to crystallize at
583 C, as shown in Fig. 2(a). With a temperature drop,
the primary (Al) phase continues to form and the liquid
composition moves to the monovariant reaction line of L !
(Al) þ (Si) along the bivariant reaction surface of L ! (Al)
because of the decrease of Al content in the liquid, as plotted
in Fig. 1(a). After the liquid composition reaches the
monovariant reaction line of L ! (Al) þ (Si), the eutectic
(Al) and (Si) phases crystallize and the liquid composition
moves to the eutectic reaction point E1 of L ! (Al) þ (Si) þ
(579 C). After liquid composition reaches point E1 , the
1310
S. Lee, B. Kim and S. Lee
Table 2 Formation characteristics of intermetallic compounds observed in four different specimens of low and high Fe-containing
eutectic Al-Si alloys solidified under slowly- and rapidly-cooled conditions.
Specimen
Alloy 1
Slowly
cooled
Primary
phase
(associated)
Primary
phase
(single)
Eutectic
phase
(single)
Morphology
Size,
l/mm
All
—
—
Needle
101 102
Minority
þ
—
—
Majority
—
—
Needle
103 104
Minority
—
—
Needle
101 102
All
—
—
Needle
100 101
Majority
—
—
Needle
102 103
Majority
—
—
Needle
100 101
Irregular (core)
Alloy 2
102 103
Needle (shell)
Alloy 2
Alloy 1
Rapidly
cooled
Frequency
Here, the relative frequencies among all intermetallic compounds except the matrix were roughly expressed. In the case of needle-shaped intermetallic, the
size is the lengthwise value.
eutectic (Al), (Si), and phases crystallize until all of the
liquid is exhausted.
For alloy 2, the first solid phase begins to crystallize at
782 C, as shown in Fig. 2(b). As the temperature drops, the
primary phase continues to precipitate out of the liquid; the
liquid composition moves to the monovariant reaction line
P1 -U1 of L ! þ along the bivariant reaction surface of
L ! , as plotted in Fig. 1(b). As the temperature drops
further, the and phases crystallize and the liquid
composition moves to the quasi-peritectic reaction point U1
of L þ ! þ (656 C) along the line P1 -U1 . At point U1 ,
the primary phase begins to crystallize. As both and phases crystallize, the liquid composition moves to the quasiperitectic reaction point U3 of L þ ! (Al) þ (612 C)
along the monovariant reaction line U1 -U3 of L ! þ (612656 C) because of the decrease of Si and Fe in the
liquid. At point U3 , the primary (Al) phase begins to
crystallize. The formations of the (Al) and phases continue
until the liquid composition reaches point E1 along the
monovariant reaction line U3 -E1 of L ! (Al) þ (579612 C). Finally, the eutectic phase, as in alloy 1, is
formed together with the (Al) and (Si) phases by the ternary
eutectic reaction E1 until all of the liquid is exhausted.
3.2 Microstructure
The formation characteristics of the intermetallic compounds observed in four specimens of low and high Fecontaining eutectic Al-Si alloys solidified under slowly- and
rapidly-cooled conditions are summarized in Table 2. The
chemical compositions of intermetallic compounds measured
by EDS are indicated in Table 3. The representative microstructures of the slowly-cooled low Fe-containing eutectic
Al-Si alloy are shown in Fig. 3. For the slowly-cooled alloy
1, the microstructure consists of eutectic (Al) and (Si) phases
of typical lamellar structure together with very few primary
(Al) phases, as shown in Fig. 3(a). A small amount of needleshaped eutectic phase was observed in Fig. 3(b). As shown
in Fig. 3(c), the eutectic phase is the only intermetallic
compound found in the microstructure of alloy 1, being
evenly distributed in all over the microstructure. The lengths
of the eutectic phase needles range from 10100 mm, as
shown in Fig. 3(d).
Table 3 Average chemical compositions of intermetallic compounds
observed in two different specimens of low and high Fe-containing
eutectic Al-Si alloys solidified under slowly-cooled conditions.
Chemical composition (mole%)
Specimen
Phase
Number of
measurements
Al
Fe
Si
Alloy 1
2
67.6
13.8
18.6
2
68.3
18.2
13.5
7
66.7
14.9
18.4
Alloy 2
Figure 4 shows the representative microstructures of the
slowly-cooled high Fe-containing eutectic Al-Si alloy. For
the slowly-cooled alloy 2, a large amount of coarse needleshaped precipitate was observed in Figs. 4(a) and 4(b). From
the results of combined analyses of SEM(BSE) and EDS,
it was verified that the coarse needle-shaped precipitate was
the phase. As shown in Fig. 4(c) and its inset, the existing
precipitates were predominantly the coarse single phase
and were rarely the associated phases of þ . The coarse
primary phase was visible to the naked eye because of the
length range of 110 mm and the thickness of several
hundred micrometers, as shown in Figs. 4(c) and 4(d). A
small amount of eutectic phase needles with a length range
of 10100 mm was also found in the eutectic regions of the
(Al) and (Si) phases.
Figure 5 shows OM images of the rapidly-cooled low
Fe-containing eutectic Al-Si alloy (alloy 1). A fine dendritic
primary (Al) phase, together with very fine eutectic (Al) and
(Si) phases, was observed in Fig. 5(a). Despite the neareutectic composition of 12.4 mass%Si, the amount of the
dendritic primary (Al) phase was much higher than expected.
As shown in Fig. 5(b), a small amount of very fine needleshaped eutectic phase with a length range of 110 mm was
present in the eutectic region.
Figures 6 and 7 show OM and SEM(BSE) images,
respectively, of the rapidly-cooled high Fe-containing eutectic Al-Si alloy (alloy 2). A large number of huge needleshaped primary phases with a length range of 1001000
mm was observed in Fig. 6(a). A very large number of tiny
needle-shaped eutectic phases with a length range of 110
mm, together with eutectic (Al) and (Si) phases, was also
observed in Fig. 6(b). As shown in Fig. 7, it was verified that
Prediction of Solidification Paths in Fe-Containing Eutectic Al-Si Alloys
Fig. 3 Representative microstructural features of slowly-cooled alloy 1; (a) low magnification optical micrograph, (b) enlarged view of
(a), (c) low magnification back-scattered scanning electron micrograph, and (d) enlarged view of (c).
Fig. 4 Representative microstructural features of slowly-cooled alloy 2; (a) low magnification optical micrograph, (b) enlarged view of
(a), (c) low magnification back-scattered scanning electron micrograph, and (d) enlarged view of (c). Inset in (c) shows the typical
example of incomplete quasi-peritectic reaction of L þ ! (Al) þ .
1311
1312
S. Lee, B. Kim and S. Lee
Fig. 5 Representative microstructural features of rapidly-cooled alloy 1; (a) low magnification optical micrograph and (b) enlarged view
of (a).
Fig. 6 Representative microstructural features of rapidly-cooled alloy 2; (a) low magnification optical micrograph and (b) enlarged view
of (a).
4.
Fig. 7 Back-scattered scanning electron micrograph of rapidly-cooled
alloy 2.
the length distribution of the existing phases was very
diverse from 1 mm to 1000 mm and divided into two similar
sized groups.
Discussion
4.1 Solidification path-dependent phase formation
As mentioned in Part I, among the twenty seven ternary
invariant reactions of the Al-Si-Fe system, only five can form
phase. For the Fe-containing eutectic Al-Si alloys selected
in this study, the formation of phase can occur by the
U1 quasi-peritectic, U3 quasi-peritectic, and/or E1 eutectic
reactions.
For the slowly-cooled alloy 1, the fine eutectic phases
were a unique type of Fe-rich intermetallic compound
observed in the final microstructure, as shown in Fig. 3(c).
The E1 ternary eutectic reaction has the lowest reaction
temperature (579 C) among the five ternary invariant
reactions that can form phase. If an initial composition,
which can form phase by the E1 ternary eutectic reaction
exclusively, is adopted in Fe-containing Al-Si alloys, the
length of the phase can be minimized because of low
starting temperature for phase formation and growth
constraint within the eutectic regions of the (Al) and (Si)
phases. The initial composition is decided by the critical Fe
content (FeCR ),7) which is plotted by the dash-dot line in
Fig. 1(a). In the case of alloy 1, the fine eutectic phase
among various Fe-rich intermetallic compounds was only
formed by the E1 ternary eutectic reaction of L ! (Al) þ
Prediction of Solidification Paths in Fe-Containing Eutectic Al-Si Alloys
(Si) þ because the initial composition was below FeCR .
Moreover, the phase was widely and evenly distributed in
all over the microstructure because the matrix of alloy 1 with
the near-eutectic composition was mostly consisted of
eutectic (Al) and (Si) phases with very few primary (Al)
phases.
For the slowly-cooled alloy 2, the most dominant
intermetallic compound in the final microstructure was the
very coarse needle-shaped primary phase, as shown in
Figs. 4(c) and 4(d). Figure 8 exhibits the first derivative of
the cooling curve for the slowly-cooled alloy 2. The peaks
identified in Fig. 8 closely correspond to the invariant
reactions of the solidification sequence, as was predicted
thermodynamically in Fig. 2(b). In particular, prominent
peaks corresponding to the formation reactions of the phase
such as the U1 quasi-peritectic reaction of L þ ! þ ,
the U3 quasi-peritectic reaction of L þ ! (Al) þ and the
E1 eutectic reaction of L ! (Al) þ (Si) þ were detected
(Fig. 8). The U1 quasi-peritectic reaction, together with the
P2 peritectic reaction in Fe-containing hypereutectic Al-Si
alloys, has the highest reaction temperature (656 C) among
the five ternary invariant reactions that can form phase
(In the strict sense, P2 peritectic reaction temperature is very
slightly higher than the U1 quasi-peritectic reaction temper-
ature.). If the initial composition, which can form phase by
the U1 quasi-peritectic reaction, is adopted in Fe-containing
Al-Si alloys, the length of the phase can be maximized. In
the case of the slowly-cooled alloy 2, the very coarse needleshaped primary phase predominantly formed because of the
high starting temperature for phase formation corresponds
to point U1 (656 C), the wide formation temperature range
corresponds to line U1 -U3 -E1 (579656 C), and the unconstraint growth in the liquid without the formation of the
primary (Al) phase during the solidification temperature
range of 612656 C. Regarding the solid phase consumption
by a quasi-peritectic reaction, as shown in Fig. 4(c) and its
inset, the absence of primary phase is the result of complete
U1 quasi-peritectic reaction and the presence of primary phase is the result of incomplete U3 quasi-peritectic reaction.
Figure 9 shows the calculated equilibrium phase fractions
of the two eutectic Al-Si alloys selected in this study. For
alloy 1, the calculated equilibrium amount of the primary
(Al) phase at the E1 eutectic reaction temperature is about
4 mass% (Fig. 9(a)). However, for the rapidly-cooled alloy 1,
the amount of the dendritic primary (Al) phase in Fig. 5(a)
was much higher than the predicted amount. This is because
the formation kinetics of eutectic (Al) and (Si) phases is
retarded by the induced high cooling rate. As shown in
Fig. 5(b), a very fine phase, together with eutectic (Al) and
(Si) phases, was observed because the initial composition
was below FeCR , as in the slowly-cooled case. It is thought
that the actual amount of the eutectic phase is smaller than
that of the equilibrium condition (Fig. 9(a)) because of the
increased formation of supersaturated (Al) phase.
For the rapidly-cooled alloy 2, a coarse straight needleshaped primary phase, which was unlike a curved needleshaped primary phase observed in Part I, was predominantly observed in the final microstructure, as shown in
Fig. 7. In this case, it is predicted that the primary phase is
mainly formed by the reaction of L ! (583649 C), as
shown in Fig. 9(b). It is noted that there is no change of
liquidus composition prior to the reaction given the absence
of the primary and phases, as identified in Fig. 7.
Therefore, it is thought that the primary phase grows easily
straight along the lengthwise direction in the liquid without
the dendritic primary (Al) phase.
Fig. 8 First derivative of the cooling curve for slowly-cooled alloy 2.
(a)
579°C
100
90
(Al)
60
50
40
30
20
0
500
656°C
E1
U1
Liquid
80
E1
70
10
579°C
100
90
Phase (mass%)
Phase (mass%)
80
(b)
Liquid
70
(Al)
60
50
40
β
30
20
(Si)
10
β
550
1313
600
650
700
750
Temperature, T / °C
Fig. 9
800
850
900
0
500
α
(Si)
550
600
650
γ
700
750
Temperature, T / °C
Calculated equilibrium phase fractions; (a) alloy 1 and (b) alloy 2.
800
850
900
1314
S. Lee, B. Kim and S. Lee
(a)
(b)
Fig. 10 Effect of Fe content on the equilibrium fraction of the phase; (a) Al-7 mass%Si and (b) Al-12.8 mass%Si alloys.
4.2 Effect of Fe content on phase formation
For the two slowly-cooled alloys, with increasing Fe
content above 8 mass%, the primary phase formed by the
U1 quasi-peritectic reaction of L þ ! þ (656 C) was
more dominant than the eutectic phase formed by the E1
eutectic reaction of L ! (Al) þ (Si) þ (579 C). For the
two rapidly-cooled alloys, with increasing Fe content,
the primary phase formed by the reaction of L ! (583649 C) was more dominant than the eutectic phase.
In the both cases of the slowly- and rapidly-cooled eutectic
Al-Si alloys, the length of the phase increased by about two
orders of magnitude. It is verified that increasing the Fe
content promotes increases in the amount and length of the phase. This is because the supply of Fe atoms for phase
growth is facilitated by increases in the phase formation
temperature (related to the diffusivity of Fe atoms) and the
solute Fe concentration in the liquid (related to the jump
frequency of Fe atoms).
4.3 Effect of cooling rate on phase formation
In both cases of alloy 1 and 2, by increasing the cooling
rate to a level of 250300 C/s, the length of the phase
decreased by about an order of magnitude. It is thought that
increasing the cooling rate inhibits the formation of the
coarse phase because of an insufficient exhaustion of
incubation time for nucleation and a shortage of time for
lengthwise growth. For alloy 2, with increasing the cooling
rate, it was found that the length distribution of the existing phases was very diverse from 1 mm to 1000 mm and divided
into two similar sized groups of several micrometers and
several hundred micrometers. The reasons for the diverse
length distribution are the high starting temperature of the
phase formation and the wide formation temperature
range. The reasons for dividing the two groups are that the
amount of the primary phase fails to reach the maximum
amount at thermodynamic equilibrium due to the induced
high cooling rate and the formation of eutectic phase is
promoted at or below the ternary eutectic temperature. It may
be argued that the amount and length of the phase in alloy
2 cannot be completely controlled by high cooling rate of
250300 C/s.
Combined effect of Fe and Si contents on phase
formation
Through comparing the results of the high Fe-containing
hypoeutectic Al-Si alloys in Part I, it was found that the
length of the phase in eutectic Al-Si alloys was about an
order of magnitude longer than that in hypoeutectic Al-Si
alloys. It means that the more Si content is close to the
eutectic composition, the more effective is the increase of the
phase length by increasing Fe content. In order to analyze
the reason, Al-7 mass%Si and Al-12.8 mass%Si alloys were
representatively selected in hypereutectic and eutectic alloys,
respectively. Thermodynamic calculations of the two selected Al-Si alloys were performed using the updated database.
Figure 10 shows the calculated equilibrium fraction, which is
the theoretical maximum amount, of the phase depending
on temperature according to Fe content in the two Al-Si
alloys. Below the U3 quasi-peritectic reaction temperature,
the equilibrium fraction of the phase according to Fe
content in the eutectic Al-Si alloys is remarkably consistent
with that in the hypereutectic Al-Si alloys. However, for the
eutectic Al-Si alloys with the high Fe level, the formation of
the phase also can occur in the high temperature range
between U1 and U3 quasi-peritectic reactions. This is because
the U1 quasi-peritectic reaction is becoming increasingly
stable with increasing Si content to the eutectic composition.
Moreover, the phase growth can proceed favorably and
quickly in the high temperature region without the dendritic
primary (Al) phase.
Figure 11 shows the calculated equilibrium temperature,
which is the theoretical formation starting temperature, of phase depending on Fe content in the two Al-Si alloys. The
Fe content that the formation starting temperature of phase
begins to increase in the eutectic Al-Si alloys is completely
corresponded to the Fe level at E1 ternary eutectic reaction
and is higher than that in the hypoeutectic Al-Si alloys. This
is because the allowed Fe content to prevent the formation
of the primary phase becomes higher with increasing Si
content (It is also similar to the concept of FeCR .). The slope
of a curve shown in Fig. 11 means the Fe content sensitivity
on the starting temperature for phase formation. The slope
of the eutectic Al-Si alloys is steeper than that of the
4.4
Prediction of Solidification Paths in Fe-Containing Eutectic Al-Si Alloys
1315
(5) For rapidly-cooled high Fe-containing eutectic Al-Si
alloys, the length distribution of the phases was very
diverse from 100 mm to 103 mm and divided into the two
similar sized groups by rapid cooling.
(6) Acceleration of the phase formation by adding Fe in
eutectic Al-Si alloys is more effective than that in
hypoeutectic Al-Si alloys.
(7) For eutectic Al-Si alloys, the length of the phase can
be minimized by decreasing Fe content below
0.67 mass%.
(8) For eutectic Al-Si alloys, the length of the phase
can be maximized by increasing Fe content above
2.8 mass%.
Fig. 11 Effect of Fe content on the equilibrium temperature of the phase
in Al-7 mass%Si and Al-12.8 mass%Si alloys.
hypereutectic Al-Si alloys. It means that Fe content in
eutectic Al-Si alloys is more effective than that in the
hypoeutectic Al-Si alloys on the formation behavior of the
primary phase.
For the eutectic Al-Si alloys, as shown in the calculated
result of Fig. 11, it is thought that the length of the phase
can be minimized when Fe content is decreased to somewhat
below the Fe level at the E1 ternary eutectic reaction
(;0:67 mass%) and maximized when Fe content is increased
to somewhat above the Fe level at the U1 ternary quasiperitectic reaction (;2:8 mass%). For the addition of Fe
content above 2.8 mass%, there is no effect of the added Fe
content on the starting temperature for phase formation.
However, the added Fe content increases the solute Fe
concentration in the liquid and results in promoting the phase growth.
5.
Conclusions
Thermodynamic calculations were carried out to predict
the solidification sequences and formation behaviors of the phase in Fe-containing eutectic Al-Si alloys. These predictions were verified by appropriate experimentation. The main
findings are as follows.
(1) For slowly-cooled high Fe-containing eutectic Al-Si
alloys, the very coarse needle-shaped primary phase
with a length range of 103 104 mm was predominantly
formed by the quasi-peritectic reaction of L þ !
þ (656 C).
(2) For rapidly-cooled high Fe-containing eutectic Al-Si
alloys, the coarse needle-shaped primary phase with a
length range of 102 103 mm was considerably formed
by the reaction of L ! (583649 C) without growth
restriction by the dendritic primary (Al) phase.
(3) By increasing the Fe content above 8 mass%, the length
of the phase increased by about two orders of
magnitude.
(4) By increasing the cooling rate to a level of 250
300 C/s, the length of the phase decreased by about
an order of magnitude.
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