Solidification structure
5.4 Aluminium alloys
5.4.3 Nucleation and growth of the solid (grain refinement)
The addition of titanium in various forms into aluminium alloys has been found to have a strong effect in nucleating the primary aluminium phase.
It is instructive to consider the way in which this happens.
Titanium in solution in the liquid metal at a
1 o4
10-3’
1 k m 10 pm 100 k m 1 mm
Inclusion diameter
10 mm Figure 5.35 An example of inclusion content in an A1 alloy.
concentration above about 0.15 weight per cent would be expected to be precipitated as TiAI, in the peritectic reaction (Figure 5.36). There is no doubt that TiA13 is an active nucleus for aluminium because TiA1, is found at the centres of aluminium grains, and there is a well-established orientation relationship between the lattices of the two phases (Davies et al. 1970).
However, there are two major problems: we have to ask the questions (i) ‘How does the TiA13 phase itself nucleate?’ and (ii) ‘Why does the TiA1, phase exist below the 0.15 per cent Ti, where it should be unstable, and should go into solution as indicated on the equilibrium phase diagram’ (Figure 5.36).
Results of several researchers shown in Figure 5.37 illustrate that the effect of titanium in the grain refinement of aluminium starts at much lower concentrations, below 0.01 titanium.
How titanium can be effective at concentrations lower than 0.15 per cent has remained a mystery in A1-Ti liquid solutions until the epoch-making research by Schumacher and Greer (1993 and 1994).
These researchers carried out their studies on an amorphous aluminium alloy, as an analogue of the liquid state. Nucleation is more easily observed since the kinetics of reaction are 1016 times slower.
Using TEM (transmission electron microscopy) they observed that TiA13 was present as adsorbed layers on TiB2 crystals, and so its existence was stabilized
1000
900
800
a
v
2 2 700
-
W
E
e
600
500
400
Liquid + TiAI,
/
/ a + TAI,
/ / / I I
I I I I I
0.5 1 .o 1.5 2.0
Titanium (wt per cent) Figure 5.36 Binary AI-Ti phase diagram.
Solidification structure I5 1
assumed by Greer and colleagues (2001) to be controlled by the rate at which solute can diffuse through the segregated region ahead to the advancing front. They use a growth restriction parameter Q defined as
107
1 o6
105
h i o 4
v
5
.
C._
k-
! c3 10'
1oi
IO'
I I I /
/ /
/ / Davies et a/ (1 970) A
r
' Youdelis and Yang (1 982) 0Delamore et a/ (1971)
I I I I I
1'04 10-3 io-* io-' 1
Ti concentration (wt per cent)
Figure 5.37 Increase in grain refinement with increasing titanium addition, especially at the peritectic 0.15 Ti.
at lower levels of Ti than would be expected from the phase diagram, and it was thereby effective in nucleating aluminium.
It is worth noting, however, that in real castings (as contrasted with laboratory experiments), the action of titanium in grain refinement may be rather different. In this case the titanium is added as a master alloy that already contains TiA1, particles in suspension. As the master alloy disperses in the bulk melt, the TiAI, particles will tend to dissolve, and sink to the bottom of the furnace. However, this takes time of the order of an hour or so, allowing plenty of time in practice for the treatment and casting of the melt. The addition of boron to the master alloy greatly increases the effectiveness of grain refinement. Although it was thought that the TiB2 might act in several ways (review: McCartney 1989), it seems most likely that the stabilization of TiAl, on its surface might be the most important factor, lengthening the life of the active TiA1, particles.
Effective grain refinement, however, seems to require more than simply the nucleation of new grains. A second important factor is the suppression of their growth. For complex systems, where many solutes are present, the rate of growth of grains is
Q = m(k - l)Co (5.22)
Although they note that this relation should be modified by the rate of diffusion D of the solute, these factors are not well known for many solutes.
They therefore assume that the rates of diffusion are fairly constant for all solutes of interest in aluminium (this is not far from the truth for the substitutional solutes shown in Figure I .6). Thus the potentially more accurate summation of the effects of different solutes in solution, weighted inversely by their diffusivities as proposed by Hodaj and Durand (1 997), is neglected in favour of a simple summation of Q values. The effect is shown in Figure 5.38. Initially, grain size clearly decreases with increasing total values of Q.
h
3
0 AI-Si-Fe-Ti (c3wt%sI) AI-Si-Ti (>3wt%Si) A o AI-Si-Ti AI-FeTi (<3wt%Si)
M AI-Ti
n
0 20 40 60 80
Z(ki - 1) miCo
Figure 5.38 Efect of the growth restriction ,fucfor on the grain size of various A1 a1loj.Y (Greer et al. 2001).
The subsequent apparent growth of grains with increasing Q above about 20 is thought to be (Greer 2002) the result of the special effect of Si in 'poisoning' the grain refinement action of Ti at Si contents over 3 per cent. The higher Q data is defined only by alloys with Si contents above 3 per cent.
For other solutes at high Q , particularly Cu, it is thought that the grain size remains small. It seems that the attainment of fine grain size in AI-Si casting alloys has fundamental limitations, the attainable sizes being 5 to 10 times larger than those in some other casting alloys and in most wrought alloys.
O t h e r nucleation and growth effects are happening during the solidification of many A1 alloys as a result of the many solutes that are present.
both intended and unintentional.
As an example of one of these, Cao and Campbell
(2000) discovered that pFe plates (A15FeSi intermetallic) in A1-Si alloys precipitated on the wetted outside surfaces of bifilms. Initially, the pFe precipitate is sufficiently thin that it can follow the folds of the bifilm. On a fracture surface the iron- rich phase can be clearly seen through the thin oxide film that represents one half of the bifilm (Figure 5.39). At this early stage it is faithfully following the undulations of the oxide film.
However, as the pFe particle thickens, the particle becomes increasingly rigid, taking on its preferred crystalline form, and so forces the film to straighten.
Finally, the bifilm is often seen as a crack aligned along the centre of the pFe particle, or along the matrix/particle interface if the pFe happened to nucleate only on one side of the bifilm (Figure 5.40).
In Figure 5.41a the bifilm on one side of the particle has been pulled away by gas precipitation
or by shrinkage forces, opening a pore on one side of the pFe plate. Because it has been common to observe an association between pores and pFe particles, it has in the past been assumed that the pFe particles blocked the movement of feed liquid along interdendritic channels, and so caused shrinkage porosity. However, this seems most unlikely, in view of the three-dimensional access routes for feed liquid, and in view of the strong probability that pores probably cannot be formed without bifilms.
5.4.4 Modification of the AI-Si eutectic by