Journal
Dirisiotr
of Malerials
Engiv7rcring.
of Materials Processing Technology 66 11997) 153-257
School
IJ/ Applied
.%icncr.
Nuyatg
Tecl7iiolctgic~al
Ut7icersil~~,
Sit7guporc
09798,
Sittguporr
Received 16 October 1995
-._
Abstract
The effects of the Ti:B ratio, the meit holding temperature and the melt mechanical agitation on the performance of aluminum
grain refinement master alloys were investigated using small ingot castings. The results confirm that the presence of dissolved Ti
in the melt is the key factor for effective grain refinement.
omogenizing the distribution of heterogeneous nuclei in the melt by
mechanical agitation can further enhance the performance of Ti-B based master alloy. When the amount of disso!ved Ti in the
melt is less than the peritectic point, the columnar
uiaxed transition in master alloy treated melt is more sensitive to heat
transfer variation rather than to the number of Ti
particles retained in the melt. Slow cooling is recommended to avoid
columnar growth. As the holding temperature increases, inoculation fade occurs. 0 1937 EIsevier Scieace §.A.
&words:
Al-T-i: AI-B; Al grain refinement; Master alloys
The grain refinement of aluminum and its allo
normally achieved by melt inoculalion with Ti and
the form of aluminum master alloys. A great deal of
experimental work has been carried out in the past in
trying to understand the factors that affect the performance of grain refinement and the possible mechanisms
involved [l-3]. Experimental results [4,5] have shown
that Ti can significantly retie aluminum grains, whilst
the presence of B will further enhance grain refining
performance
[6]. It has been reported that the Ti/B
ratio of the master alloy is an important factor, the best
grain refining effects being obtained using master alloy
with a Ti:B ratio of 51 in weight percentage [7].
Two models [4,8,9] have been proposed to explain
the mechanisms of effective grain refinement achieved
through the addition of Al-Ti-B
master alloy. One is
based on heterogeneous nucleation at insoluble TiB2
sites, in which the excess of Ti in the melt will restrict
crystal growth. The other model suggests that Ti in
* Corresponding author. Tel.: + 65 7991142; fax: + 65 7926559;
e-mail: [email protected]
0924-0136/97/$17.00 6 1997 Elsevier Science S.A. All rights reserved.
PII SO924-0136(96)02536-S
solution segregates to the TiB,/melt interface, leading
to the formation of a stabilized layer of atoms on the
TiB2 surface. The layer, a solid solution of Ti in Al, will
promote growth of Al from the melt during solidification. These models, however, are based on hypotheses
which are difficult to verify experimentally. (Al,Ti)B,
intermetallics may also serve as heterogeneous nuclei
but their stability in Al melt is still in doubt [lo]. Recent
experimental findings [l l] indicate that TiB2 particles
alone do not promote heterogeneous nucleation, whilst
precipitation of an TiAl, layer at the TiBz surface will
nucleate Al.
For Al-Ti master alloy grain refinement, a model
[ 121 attributes grain refinement to the peritectic reaction
involving TiA& particles: Liquid Al + TiAl, = solid solution Al. The newly forr .ed solid solution particles will
serve as nuclei for grains. Experimental evidence clearly
support this model [13,14].
Much of the research effort have been focused on the
effect of Ti/B ratio and the presence of other alloying
elements on the performance of a master alloy, whilst
less attention has been given to the possible effects of
other processing parameters such as the melt holding
temperature, melt agitation and the iooling rate on the
perfiormance of the grain refinement.
H. Li er al. /Jorrrnal of’Materials Proressing Tehology
254
Table 1
Chemical compositions
of commercially pure Al and the Ti-AI and B-Al master alloys (wt.‘%)
Element
Pure Al
Ti-Al
B-Al
-
Fe
Si
0.110
0.058
0.2
0.20
0.30
66 (1997) 253-257
Mg
Ga
V
Ti
0.001
0.014
0.01 I
the holding temthe cooling rate
alloys is investiand B contents,
2. Experimental
All experiments were conducted on 100 or 200 g
commercially pure aluminum ingots cast in graphite
crucibles. A master alloy of nominal composition 10
wt.%
Ti which contained flaky TiAl, particles was
used to add Ti whilst a 5 wt.% B master alloy was
chosen for the B addition. The composition of the
commercially pure Al and the B-Al and Ti-Al master alloys is given in Table 1. Note that this procedure is different to reported
procedures
in the
literature where Tibor master alloys are added which
contain both Ti and B in one master alloy, preadded, in pre-determined amounts.
Experiments were carried out to examine the effects
of four processing parameters, namely the Ti/B ratio,
superheating, the cooling rate and stirring, on the solidification structure of the ingots.
Four 100 g ingots were cast with different Ti/B
ratios, an aluminum ingot without inoculation treatment also being produced as a reference. The proportion of master alloys used in the investigation is given
in Table 2. Aluminum melts were superheated
to
725°C in the crucible by a resistance furnace, and
then pre-weighed master alloys were added. The melts
were stirred with an alumina rod at intervals of 15
Table 2
Proportion of Ti and B used in the 100 g i -got castings
Sample
Wt.% of Ti
Wt.‘% of B
Atomic ratio Ti/B
1
2
3
4
0.150
0.125
0.050
0
0
0.005
0.020
0.030
6/l (whole number)
l/2 (whole number)
-
Misc.
Al
-.
IOk 1.0
0.05
In the pre;ent work, the effects of
perature, mechanical agitation and
on the performance of the master
gated in addition to those of the Ti
through small ingot casting.
I3
0.03
5 + 0.5
0.025
0.03
0.03
Bal.
Bal.
Bai.
min for 1 h to ensure dissolution of the master alloy,
and then allowed to solidify in still air for subsequent
metallographic examination.
Five 100 g ingots with a fixed amount of master
alloys additions (0.125 wt.% Ti and 0.005 wt.% B;
Ti/B atornic ratio 6:l) were also cast at different
holding temperatures of 725, 800, 850, 900 and 95O’C
respectively. The alloy addition, holding time and
casting procedures were the same as described above.
Two 200 g ingots with the same 0.125 wt.% Ti and
0.005% B addition (Ti/B atomic ratio 6: 1) were cast
with and without additional mechanical stirring just
before the solidification
of the ingot. All other
parameters were as described earlier, except for the
stirring operation. One melt was taken from the furnace at 725°C and allowed to solidify in air without
any agitation, whilst the other melt was taken out
from the furnace at the same temperature and immediately stirred before solidification. The cooling curve
of the melts were determined using a K type thermocouple and a chart recorder. The stirring was carried
out using an alumna rod and lasted for 15 s. Stirring
stopped when the melt reached 700°C, the melt then
being allowed to cool down to the melting point and
solidify in air.
The ingots were sectioned through their centre
lines, which later were then polished and etched to
reveal their grain structures. Metallographic specimens
were taken from the centre of each ingot at the crosssection for microscopic studies. Poultant’s regent containing 30 ml HCl, 15 ml of HN03, 2.5 ml of 40%
HF and 2.5 ml of distilled water was used for
macrostructure
etching. Photomacrographs
and photomicrographs were taken at the cross sections of ingots using a 35 mm camera with a macro-zoom lens,
and a stereo microscope.
Grain size measurements were carried out on the
etched specimens. The linear-intercept
method was
used in which a test line of length L was drawn
randomly on the sample in different directions. The
length of the line varied, depending on the grain size
of each ingot. The number of grains intercepted was
counted. For each sample, 6 test lines at different
directions were drawn for grain size measurement.
‘55
Grain
Size
of
Pure
=
0.04~
cm
z
UL
65
90
115
140
Superheatin?
0
Tl
0.03
Fig. I.
ingots.
0.025
0.025
0.05
0.02
Cl 075
0.015
Effect of the Ti and B contents
on
0.1
0.01
the
average
0.005
0
B
Fig. 2. Effect of the melt
grain size of Al ingots.
190
215
Temperature
165
I
superheating
temperature
240
265
Degree
on
290
C
the avemge
grain size of Ail
The effect of the Ti and
contents on the average
grain size of the ingots is summarized in Fig. 1. The
x-axis in this figure shows the Ti and B contents in
ith the Ti content increasing to
weight percentages,
content decreases in this dire
the right whilst the
0.05 wt.% Ti and 0.02 wt.%
tion. The compositi
corresponds
to 8 Ti:
atomic ratio of 1:2 approximately, which is the stoichiometric ratio for the compound TiB?. Therefore, compositions to the right of
this ratio will have Ti contents in excess of what is
required to compound with available B. This excess Ti
will either be in sokion in the aluminium or as TM,.
By a similar reasonkg, compositions to the left of the
1:2 ratio of Ti:B will have excess 5. Again, this excess
B will be either in the form of solution in aluminium or
form the compound AlB,. The composition 0.125 wt.%
Ti and 0.005 wt.% B corresponds to a Ti:B atomic ratio
of 6:1 approximately, i.e. it has excess Ti.
Although B addition alone did reduce the average
grain size of the ingot, significant reductions occurred
for compositions where the atomic ratio of Ti:B was
increased beyond 1:2. This means that excess Ti addition significantly refined the grain structure. Even pure
Ti addition gave a substantial reduction in grain size,
albeit at high levels of addition (peritectic point).
The composition where the Ti:B atomic ratio was 6:l
appeared to give a lower grain size than the melt which
had 0. I5 wt.% Ti only (without any B). This may imply
that small additions of B are beneficial and even that
the Ti level could be replaced by small amounts of B.
For example, the alloy of composition 0.125 wt.% Ti
and 0.005 wt.% B (Ti:B = 6:l) gave an average grain
sizf: slightly smaller than the alloy which had 0.15 wt.%
Ti a,Id 0 wt.% B. Thus, the Ti level could be reduced to
0.125 wt.% from 0.15 wt.% (a saving of 17% in the
weight of Ti), without any adverse effects on the grain
size, if 0.005 wt.‘%)
should be added. This could
translate into cost savings in grain refinement. when Ti
ad&ions
could be partly traded off for smaller
amounts of
ore experimental work is required to
determine the optimum amounts of Ti and B for any
given alloy.
The plot of the melt holding temperature after master
alloy addition against the average grain size of ingots is
shown in Fig. 2. It is clear that as the holding temperature increased from 725 to 910°C, the average grain
sizes of the ingots increased linearly.
The melt temperature is a significant facnctorin grain
refinement. If the holding temperature is too high after
inoculation, some fading occurs. This can be attributed
to coarsening and settling of TiB, particles, leaving a
melt depleted of nucleating particles for efficient grain
refinement. It is known that the particle size of the TiB,
that is formed in-situ in the aluminium melt depends on
the melt temperature,
at high temperatures coarser
particles being formed that can settle tothe bottom of
the melt by virtue of their greater density.
3.3. Tile effect of mechur~ical ugitntion
Fig. 3 shows macrographs of ingots obtained with
and without mechanical agitation after the inoculation
treatment, the enhanced refinement obtained by me-
Fig.
3. Cross-sections of Ai ingots with and without
agitation.
mechanka!
256
H. Li et al. /Journal of Materials ProcessitrgTechnolog,v66 (1997) 253-257
chanical agitation being clear. For both samples,
columnar grains were observed at the top surface of the
ingots that were followed by Gner, equiaxed grains
progressively downwards to;\:ards the bottom of the
ingots. However, the average grain size in the agitated
ingot is clearly finer and more uniform across the
section, when compared to the unagitated ingot. Thus,
agitation before solidification gives a finer and more
uniformly distributed grain structure across the whole
ingot. However, agitation does not have a significant
effect on the columnar-equiaxed grain transition.
4. Diiussion
4.1. The grain refinement mechanism
The essence of grain refinement inoculation is to
promote the columnar-to-equiaxed transition during
solidification of a casting. Several models have been
proposed for this transition, of which Hunt’s steady
state model [15] gives the best insight to the transition
in a simple formulation. Three parameters are recognised by Hunt’s model as influencing this transition:
alloy composition, nucleation undercooling and nucleation site density. In the present studies, the alloy
composition has remained unchanged except for the
different levels of Ti and B and hence this variable does
not exist. The solidification rates also were essentially
constant, particularly through the solidification temperature range, and hence the nucleation undercooling
during solidification could only be influenced by the
inoculation.
Results in Fig. 1 reveal that when the atomic Ti/B
ratio is kept at 1:2 or less in the ingot, the effectiveness
of grain refinement is poor. As a matter of fact when
the atomic ratio Ti/B is equal to or less than 1:2, all of
the Ti will be consumed in forming stable TiB, which
has a large enthalpy of formation of 326.41 kJ mol - l.
In this case no dissolved Ti is left in the melt, resulting
in very limited grain refinement. The addition of 0.15
wt.% Ti alone can significantly reduce the average grain
size of the ingot. The presence of a small amount of B
can achieve even better grain refinement. In other
words, retaining a small amount of dissolved Ti in the
aluminum melt is the key requirement for effective
grain refinement.
Both ingots shown in Fig. 3 contain combined 0.13
wt.% Ti and B. The atomic Ti/B ratio is 6:1, indicating
that both TiB2 intermetallic particles and dissolved Ti
are present in the melt during ingot solidification. TiB,,
as a intermetallic phase in the aluminum melt, has a
density of 4.48 g cm- 3, which is significantly greater
than that of aluminum at 2.70 g cmd3. During the
holding time, TiB2 particles will tend to move downwards and settle at the edge and bottom of the ingot,
Fig. 4. Enlarged view of the cross-sections of Al ingots with and
without mechanical agitation, showing the columnar zones at the top
of both ingots.
leaving less particles suspended in the melt. The downward rate of these particles is a function of the particle
diameter and the viscosity of the melt, described by
Stokes’s Law. Upon solidification, a large amount of
TiB2 particles accumulated at the edge and the bottom
of the melt produced a lot of heterogeneous nucleation
sites, resulting in finest grains. Since less TiBz particles
were present at the centre of the melt during solidilication, the average grain size is larger there. Mechanical
agitation redistributes the TiB2 particles and provides a
large number of particles throughout the whole melt, as
a result an overall better grain refinement being
achieved. The results in Fig. 3 show clearly that heterogeneous nucleation played a major role in grain refinement. A greater melt holding temperature decreases the
viscosity of the melt and reduces the number of heterogeneous nucleation sites suspended in the melt, resulting in the inoculation fade shown in Fig. 3.
The appearance of columnar grains at the top of
each in.got can be attributed to enhanced heat transfer
dur!ng ingot solidification. When ingots start to cool
down and solidify, the heat removal rate will be greatest at the open top of the melt where the liquid metal is
exposed directly to room-temperature air. The relatively
high heat removal rate at the top surface of the melt
will in turn increase the temperature gradient in the
melt, promoting columnar grain growth. Mechanical
agitation that redistributes TiBz particles, however, will
not affect Ti dissolved in the melt. An enlarged view of
the ingots is shown in Fig. 4. It is evident that mechanical agitation only slightly reduced the grain size in the
ingot top columnar zone, and did little in promoting
columnar-equiaxed
transition. This result indicates
strongly that the change in the number of TiBz particles
in front of the advancing columnar-melt interface does
not have any signifirant influence on columnarequiaxed transition as long as the amount of dissolved
Ti remains the same: this is because the total amount of
Ti addition to the melt was 0.125 wt.%, which is less
than the peritectic point shown in Fig. 5. After the
formation of TiB,, the dissolved amount of Ti was
direct contact with hot emcible
waI1, resulting
in a
erature gradient in the melt during solidificaCc)). Equiaxed grain nucleation will start
is not severe. The
zones ia inoculation-
olved Ti level is lower
Fig. 5. The aluminium-rich
phase diagram.
end of the Al--Ti binary equilibrium
further reduced.
key requirement
the advancing soli
interface is AT, < AT,, where
AT, and AT, are
rees of heterogene
nucleation supercoolin
?) and
growing dendrite tip supercooling, respectively. In this
experiment,
two factors may make the columnarequiaxed transition difficult. Firstly, upon solidification,
the liquid in front of the advancing solid-liquid interface will be further depleted of Ti (or enriched wit
as shown schematically in Fig. 6(a). Fast cooling results
in a greater rate of advance of the interface, producing
severe Ti depletion. Without sufficient dissolved Ti in
the liquid, the critical AT, value required for heterogeneous nucleation based on TiB, particles can be greater.
Secondly, if the temperature gradient in the liquid at
the interface is steep, the degree of constitutional supercooling will be low, which again reduces the possibility
of early columnar-equiaxed
transition as shown in Fig.
6(b).
Consistent
with the above discussion, the fine
equiaxed grains formed along the edges of ingots are
due to the low heat removai rate where melt was in
favours a slow cooling condition.
e AI-E
and Almaste: alloys are
added, the presence of dissolved Ti in the aluminum
melt is the key factor to the achieving of effective grain
refinement.
ation is the major mechanism
mogenizing the distribution of
heterogeneous nuclei in the melt by mechamcal agitation can further enhance the performance of Tibased master alloy.
3. When the amount of dissolved Ti in the melt is less
than the peritectic point, the columnar-equiaxed
transition in master-alloy-treated
melt is more sensitive to
heat-transfer variation than it is to the number of TiB,
particles retained in the melt. Slow cooling is recommended to avoid columnar growth.
[I] L. Arnerg,
i-6.
[2] L. Arnerg,
I-13.
[3] L. Amerg,
i4--17.
[41 I. Maxwell
[51 I. Maxwell
161 A. Cibula,
171 J. Pearson
181 G.P. Jones
L. Bad;ierud and H. Kiang, Metals Tech., (Jan. 1982)
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L. Backerud and H. Kiang, Metals Tech., (Jan. 1982)
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and M.E.J. Birth, J. Metals. (Nov. 1979) 27-31.
ard J. Pearson. Metall. Trans., 7B (6) (1976) 223.-
234.
Fig. 6. Schematic representations of the advancing solid-liquid interface. (a) Ti depletion in the liquid at the interface; (b) constitutional
supercooling on fast cooling; (c) columnar-equiaxed
transition on
slow cooling.
191 G.P. Jones, NPL Rcorr, DMA(A) 19, National Physical Lab.,
Teddington, London, Nov., 1989.
[lOI D.G. McCartney, hr. Maarer. Rec., 34 (5) (1989) 247-260.
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