Metallurgy of Casting

METALLURGY OF CASTING
Prof. B. S. Murty
Department of Metallurgical and Materials Engineering
IIT Madras
Thermodynam!" and #net!" of So$df!aton:
The higher entropy of liquid in comparison to the solid causes the free energy of the liquid to
decrease at a higher rate than the solid with increase in temperature. This leads to a situation
of lower free energy for the liquid above the melting point and a lower free energy for the
solid below the melting point. This gives the driving force for the solidification of a liquid
below its melting point.
Solidification and Grain Size Strengthening
• Stages of Solidification
– Nucleation: occurs when a small piece of solid forms in the liquid
and must attain a minimum critical size before it is stable
– Growth: occurs as atoms from the liquid are attached to the tiny
solid until no liquid remains
– Both conditions are met when the free energy of the particular
phase is lower
F%. & Driving force for the solidification below the melting point
However this driving force is opposed by the increase in energy due to the creation of a new
solid!liquid interface during solidification. In fact this is the reason for the usually observed
"in" in the cooling curve during solidification which is termed as undercooling.
Cooling Curves
F%. ' #ooling curves during solidification
$ssuming that solid phase nucleates as spherical %clusters& of radius r it can be shown that
the net 'e(cess) free energy change for a single nucleus ∆G
'r)
is given by:
where γ
*+
is the solid,liquid interfacial energy.
F%. ( -ucleation barrier during solidification
*olving the above equation for the overall free energy minimi.ation one can get the critical
radius r/ 'defined as the radius at which ∆G
'r)
is ma(imum) for the formation of the solid
nucleus which can further grow to complete the solidification.
The associated energy barrier to nucleation ∆G/ is found by substituting r/ into the overall
free energy e(pression.
It s important to note the temperature!dependence of these terms i.e. that r/ ∝ 0,∆T and ∆G/
∝ 0,∆T
1
. This is the basis for obtaining fine grain microstructure on fast cooling. 2ast cooling
leads to higher undercooling which gives raise to finer nuclei and high nucleation rate
causing the formation of large number of fine grains.
Ca"t M!ro"tru!ture)
During casting solidification generally starts on the mould surface leading to the formation
of distinct microstructural domains such as chilled crystals columnar grains and finally the
coarse equia(ed grain structure at the centre of the casting as shown in 2ig. 3.
Casting or Ingot Structure
F%. * #ast microstructure
Different types of microstructures can develop in the casting depending on the temperature
gradient ahead of the liquid!solid interface and the growth velocity of the interface. These are
classified as planar cellular and dendritic interfaces as shown in 2ig. 4. $t very slow growth
rate and high gradients planar interface is stabili.ed '2ig. 5). The formation of dendritic
microstructure has to be suppressed particularly in directional solidification and .one
refining where compositional segregation is intended to be minimi.ed. $ dendritic
microstructure can lead to microsegregation and microporosity.
Growth Interfaces
• Growth of Interfaces depends on
– Concentration Gradient
– emperature Gradient
– undercooling
– Growth !ate
• ypes of Interfaces
– "lanar
– Cellular
– #endritic
– $quia%ed
F%. + Types of microstructures developed during soldification
Schematic Illustration of
Solidification &orphologies
F%. , Effect of temperature gradient 6 and growth velocity 7 on the microstrctures
So$df!aton of a "n%$e -ha"e A$$oy
$ single phase alloy as shown in 2ig. 8 solidifies usually by the nucleation and growth of
solid solution in the liquid. The cooling curve for such a solidification process is shown in
2ig. 9. The nucleation and growth of dendrites in such an alloy is shown in 2ig. :.
F%. . *olidification of a solid solution

Solidification Simulation
'Growth of Solutal #endrites: ( Cellular (utomation &odel and Its )uantitative Capabilities*+ ,- Beltran.Sanchez
and #-&- Stefanescu+ Metallurgical and Materials Transactions A+ /olume 01(+ 2eb- 3440+ p- 056.073
F%. / -ucleation ad growth of dendrites during solidification of a single phase alloy
The free.ing range which is the temperature difference between the liquidus and solidus
controls to a large e(tent formation of mushy .one 'mi(ture of liquid and solid) during
solidification which leads to the microsegregation and microshrin"age,microporosity. +ong
free.ing range alloys are those alloys which have large difference between the liquidus and
solidus causing e(tensive microsegregation,microporosity as shown in 2ig. 0;. In addition
fast cooling 'non!equilibrium cooling) can also lead to another segregation pattern called
coring as shown in 2ig. 00.
$quilibrium Solidification of a
Solid Solution (lloy
Cooling Curve
F%. 0 #ooling curve for the solidification of
single phase alloy
F%. &1 Effect of free.ing range on shrin"age
Gran Refnement
It is well "nown that metals and alloys usually solidify with coarse columnar grain
structure under normal casting conditions unless the mode of solidification is carefully
controlled. It is possible to develop fine equia(ed grains in as cast structure either by raising
the number of nucleation sites or by grain multiplication. 6rain refinement has been defined
as deliberate suppression of columnar grain growth in the ingots and castings and formation
of fine equia(ed solidification structure throughout the material. 2ig. 01 shows an e(ample of
conversion of coarse columnar grain structure to fine equia(ed structure by the addition of
;.;0<Ti in the form of $l!4Ti!0= master alloy as a grain refiner to $l. 2ine grains can be
achieved by fast cooling that ensures a high degree of undercooling. $lternatively raising the
number of nuclei by heterogeneous nucleation may also result in fine equia(ed grain structure
upon solidification. There are also other methods such as melt agitation and vibration during
solidification use of a mould coating which result in the formation of fine
columnar,equia(ed grain structure in the castings. $mong all these techniques grain
refinement by inoculation has become the most popular one due to its simplicity.
Nonequilibrium Solidification of
Solid Solution (lloys
$ffect of 2reezing !ange on Shrin8age
F%.&& #oring in single phase alloy during
non!equilibrium solidification
F%. &'>hotomacrograph of commercial
purity $l grain refined without and with
;.1< of $l!4Ti!0= master alloy
';.;0<Ti) showing coarse columnar and
fine equia(ed grains respectively
The interest in grain refinement stems from the fact that the mechanical properties of
any metal or alloy component are greatly enhanced with fine grain si.e. It is now proved
beyond doubt that the fine equia(ed grain structure imparts high yield strength high
toughness good e(trudability and uniform distribution of second phase and microporosity on
a fine scale resulting in improved machinability good surface finish resistance to hot tearing
and various other desirable properties.
?hen the grain refiner is added to the melt it ta"es some time for the nucleating sites to
show their full potency 'i.e. the grain si.e reaches a minimum which is commonly referred
to as ultimate grain size). The time required to reach ultimate grain si.e is normally referred
to as optimum contact time. $ grain refiner is called fast acting if the optimum contact time is
very short and is called slow acting otherwise. $n ideal grain refiner would be that which is
not only fast acting but also shows no or insignificant fading till long holding time 'long
lasting). 2ig. 0@ shows the photomacrographs of aluminium grain refined with such an ideal
grain refiner which gives fine equia(ed grains within 1min and whose potency is not lost
even after 01;min of holding the melt after the addition of grain refiner.
Grain refining efficiency can be defined as the efficiency of the grain refiner in
converting the coarse columnar grain structure to a fine equia(ed one. Hence while
comparing the grain refining efficiency of any two grain refiners the one that gives lower
grain si.e of $l for the same addition level can be considered as having higher grain refining
efficiency. Holding the melt beyond the optimum contact time usually leads to an increase in
the grain si.e which is referred to as fading. 2ading could be either due to the dissolution of
the nucleating sites and,or their settling or floating due to their higher or lower density
respectively with respect to the melt. ?hen an $l alloy containing certain alloying elements
such as #r Ar +i and high amounts of *i is grain refined with a conventional grain refiner
such as $l!4Ti!0= master alloy at the usual addition level of ;.;0<Ti only partial grain
refinement or a complete absence of grain refinement is observed which is not the case in the
absence of such elements. This is commonly referred to as poisoning effect.
$ number of methods of production of the grain refiners '$l!Ti!= master alloys) have
been reported in the literature such as the reaction of fluoride salts with $l reaction of
F%.&( >hotomacrographs of $l
grain refined with an $l!4Ti!0=
master alloy 'numbers on the
macrographs represent the holding
time in min).
molten aluminium with o(ides

thermal reduction and electrolysis melting of elemental
blends mechanical alloying and rapid solidification processing. $mong the above techniques
the one involving the reaction of halide salts such as B
1
Ti2
5
and B=2
3
with molten $l has
become quite popular. $ clean melt with good control on the alloy composition is usually
achieved by this technique. However the nature of reaction between the salts and molten $l
was not "nown till recently e(cepting that it is highly e(othermic.Cecent detailed studies
have shown that the overall reaction could be e(pressed as follows.
01D$lE F @GB
1
Ti2
5
H F 1GB=2
3
H ⇒ 1GTi$l
@
H F GTi=
1
H F 4GB$l2
3
H F B
@
$l2
5

$ttempts have been made in the past to achieve fine equia(ed grain structure in as!cast
$l alloys by small additions of a number of elements li"e Ti = Ar -b 7 ? Ta #e etc.
'commonly referred to as hardeners) to molten metal prior to casting.

$mong these elements
Ti is the most efficient grain refiner. =oron addition is "nown to improve the grain refining
efficiency of Ti. In the early stages of development of this field Ti and = were introduced
into the melt as reducible halide salts such as B
1
Ti2
5
and B=2
3
. However this practice was
later abandoned due to the difficulties in controlling the Ti and = contents high dross and
large fumes. Ti and = additions to $l in the form of master alloys li"e $l!Ti and $l!Ti!=
were found to be more effective than in the salts. In addition these master alloys offer better
cleanliness and improved product quality. 6rain refinement by $l!Ti!= master alloys is
believed to be due to the heterogeneous nucleating sites such as Ti$l
@
and Ti=
1
particles
present in these alloys.
It is in general observed that ternary $l!Ti!= master alloys have better grain refining
efficiency than the binary $l!Ti master alloys. 2ig. 03 shows clearly that even with a ;.;4<Ti
addition in the form of $l!4Ti master alloy the grain si.e of $l is larger than that achieved
with the conventional addition level ';.;0<Ti) of $l!4Ti!0= master alloy. In addition $l!
4Ti!0= usually has a better grain refining efficiency than $l!@Ti!0= as shown in 2ig. 04. The
higher grain refining efficiency of $l!4Ti!0= master alloy compared to $l!@Ti!0= could be
due to the larger fraction of nucleating sites available in former case. In case of $l alloys
containing poisoning elements such as +i Ar #r and *i $l!@Ti!0= and high = containing
master alloys appear to be better grain refiners than an $l!4Ti!0= master alloy.

F%. &*6rain si.e of commercial purity
$l as a function of Ti addition level
with binary $l!Ti master alloy at
different holding times
F%. &+#omparison of grain refining
efficiencies of $l!4Ti!0= '4,0Ti=$l)
$l!@Ti!0= '@,0Ti=$l) and $l!4Ti master
alloys for the grain refinement of $l.
$ number of test procedures such as $lcan test $lcoa test the B=I ring test and
$luminium $ssociation test e(ist for assessing the grain refining efficiency of a grain refiner.
However most of these tests are associated with high cooling rates that may influence the
assessment of performance of the grain refiners. Thermal analysis technique has also been
used to assess the performance of grain refiners in case of a number of foundry alloys. This
technique is quite fast and is being used by many foundries.
$ number of manufacturing process parameters such as Ti,= ratio reaction temperature
reaction time sequence of addition of salts and stirring conditions strongly influence the grain
refining efficiency of the $l!Ti!= master alloys prepared by the salt route. The literature and
e(perience of the authors have shown that the microstructure of the master alloy particularly
the morphology si.e and si.e distribution of Ti$l
@
particles has a strong bearing on the grain
refining characteristics of the master alloy. =loc"y Ti$l
@
particles are "nown to ma"e the
master alloy fast acting while plate li"e particles result in a slow acting one. 2igure 5a and b
show the bloc"y and plate li"e Ti$l
@
particles respectively in $l!Ti!= master alloys. $ higher
reaction temperature 'I0;;;
;
#) during master alloy preparation is "nown to result in plate
li"e Ti$l
@
particles while at lower temperatures 'I9;;
;
#) bloc"y ones are formed. The
present authors have also shown that thermo!mechanical treatments of the master alloys lead
to improved grain refining efficiency. The alloying elements have a strong role to play some
of them such as #u Mg An etc aiding in grain refinement while some others such as #r Ar
+i and *i causing poisoning.
The mechanism of grain refinement in $l after the addition of $l!Ti!= master alloys is
still a subJect of controversy. *everal mechanisms have been postulated but no clear
consensus has emerged as yet. The mechanisms available can be classified as nucleant
paradigms and solute paradigms. The nucleant paradigms deal with the heterogeneous
nucleation of solid $l on some nucleating sites. The solute paradigms incorporate the
influence of solute elements on the grain refinement process. The theories available so far on
the mechanism of grain refinement can be grouped as '0) carbide,boride theory '1) phase
diagram,peritectic theory '@) peritectic hul" theory '3) hypernucleation theory '4) duple(
nucleation theory and '5) solute theory. However neither of these theories can provide a
detailed reasoning for all the e(perimental observations. There is almost certainly more than
one mechanism operating during grain refinement which in turn may depend on the nucleant
or grain refiner used the alloy being cast and the processing conditions employed. 2urther a
more systematic wor" on this subJect is necessary to formulate a universally acceptable set of
grain refinement mechanisms.
Modf!aton of A$2S a$$oy me$t"
The eutectic mi(ture in $l!*i alloys contains *i needles which deteriorate the
mechanical properties of these alloys. The eutectic *i responds favourably to modification by
fast cooling as well as by some chemical modifier such as -a and *r that changes the
morphology of eutectic *i from acicular to fine fibrous one. 2ig. 05 gives an e(ample of the
effect of modification on the microstructure. It has been confirmed by optical metallography
that the addition of small amounts of sodium containing flu(es caused the eutectic *i to
solidify with a fine apparently globular morphology whereas untreated alloy contains the *i
phase in the form of large plates with sharp sides and ends 'acicular *i). ?ith high cooling
rate as in the case of chill casting the fla"y nature of the eutectic *i changes to fine fibrous
morphology. =ut the growth mechanism of *i in the chemically modified alloy is quite
different from that in the chill cast alloy.
F%. &, Effect of modification on the microstructure
The *i platelets in an unmodified alloy seem to be disconnected when viewed at low
magnification and without deep etching. #onsequently it was believed that each of the coarse
*i particles is an isolated crystal. +ater it was observed that radial *i particles grew across the
primary $l dendrites and concluded to be grown in three dimensions. The fla"y nature of the
unmodified *i was later confirmed by electron microscopy. $t relatively fast cooling rates
such as in chill casting the $l!*i eutectic is much finer and the *i assumes fibrous
morphology.
*everal elements have been found to produce fibrous eutectic *i structure that include
*odium *trontium >otassium Cubedium #esium #alcium =arium +anthanum and
Ktterbium. $lso the additions of elements such as $ntimony *elenium and #admium have
found to reduce the coarse lamellar *i to fine acicular structure. $mong all the above
elements -a in the range of ;.;;4!;.;0< is considered to be the most effective modifier. *r
in the range ;.;1!;.;3< is widely used for modification in the industries. =ecause of the to(ic
nature of *b it is being used only by primary industries.
The addition of modifier not only changes the *i morphology from platelets to fibrous
forms but it also shift the eutectic composition towards higher *i content. +ater it has been
reported that in addition to shifting the eutectic composition to a higher *i level chemical
modification lowered the eutectic free.ing temperature from 488 ! 453
o
#.
The selection of modifier depends upon several factors such as its modification
efficiency which is related to the number of twins produced in modified *i fibres. The twin
spacing measured with different modifiers indicate minimum twin spacing or greatest number
of twins in the case of -a modification which clearly indicates that -a is a better modifier
than other elements. However the high reactivity of -a ma"es it difficult to handle. $lso
since the solubility of -a in $l is very poor ';.;0<) manufacture of $l base master!alloy is
not preferred. The high vapour pressure of -a ';.1 atm at 8@;
o
#) causes the loss of large
amount of -a and hence the control of -a retention in the melt is quite difficult.
Ceview of literature shows that a great deal of effort has been made to replace -a by
*r for modification of $l!*i alloys. Important features of *r is that it has a low o(idation
sensitivity loss by vapori.ation is very low compared to -a and hence better control in the
melt solid solubility in $l is much higher than -a and as such $l base master alloys can be
prepared for addition as modifiers. These have led towards a growing importance of *r as a
modifier and master alloys containing different amounts of *r has been developed for
industrial uses. The *r modification not only alters the *i morphology but it changes the α! $l
percentage as well as its morphology in the case of eutectic $l!*i alloy. It has been observed
that increasing the amount of *r from ;.;04!;.;@8< promotes the growth of columnar
dendrites and results in a remar"able increase in the amount of dendritic α!phase by around
1@5< compared to the unmodified alloy. The increase in the amount of α!phase can be
e(plained in terms of the shift in the eutectic to higher *i level and the lowering of the
eutectic temperature.
In hypoeutectic $l!*i alloys there are three different possible eutectic nucleation and
growth modes depending upon the solidification conditions. These are the nucleation and
growth of eutectic 'a) from the mould wall 'b) primary $l dendrites and 'c) heterogeneous
nucleants in the interdendritic liquid. The *r addition has found to alter the growth modes.It
has been found that *r addition increases the volume percentage of the porosity pore si.e and
pore number in different casting configuration. The pore morphology changes from irregular
interconnected in appearance to more rounded and distorted It has been reported that *r has
been found to increase the amount of surface defects in commercial alloys. $lloys of near
eutectic composition are also reported to be prone to surface defects when modified with *r.
#ontrary to the above findings there are reports which indicate that *r modification increases
the total volume of the porosity. *ince the hydrogen content is not altered by the addition of
*r it is argued that the porosity will not be changed. However the porosity is found to be
widely distributed in the casting with the addition of modifier such as *r. It has been observed
that the total shrin"age which is an alloy property is not at all affected by the modification.
However distribution of shrin"age between the micro piping and the micro shrin"age has
been noticed. ?ith the modification some of the piping is reduced and the amount of this
reduction appears as distributed porosity within the castings. Thus the casting appears as more
porous not due to more dissolved gas but due to the redistribution of porosity from primary
piping to micro porosity.
The effect of *b addition as a modifier has been studied in detail and found that unli"e
with alloys modified by -a or *r the properties of $l!*i alloys treated with *b are not
dependant on processing variables such as time remelting and degassing. Though the *b
treatment does not produce eutectic *ilicon with fibrous morphology as in the case of -a or
*r considerable reduction in si.e as well as shape is obtained. The under cooling produced by
the addition of small amounts of *b to eutectic $l!*i alloy showed the presence of α!$l in the
microstructure. The high temperature LCD studies carried out on eutectic $l!*i alloys
revealed the segregation of *i atom with the addition of *b which in turn promotes the
nucleation of *i atom aggregates during the solidification thus modifying the shape and si.e
of the *i phase in the solid state.
+imited data is available in the literature on the modification effect of al"ali al"aline
earth metals and rare earth metals. However the studies carried out on the effect of =a #a K
and Kb as a modifier in hypoeutectic alloys revealed that these elements modify the eutectic
*i in different degrees. $mong the four =a gave the best result the addition of which of the
order of :;; ppm gave fully modified fibrous eutectic *i. #a was also found to be a good
modifier to get fibrous *i morphology and the modification level was found to increase with
increase of #a addition. K and Kb addition have been found to modify the eutectic *i to fine
plate!li"e structure. The additions of these elements also have been found to promote changes
in the nucleation and growth mode of the eutectic. 2urther #a and K cause eutectic growth to
occur with a strong dependency on the thermal gradient i.e. eutectic growth from the surface
towards the centre on a macro scale.
There are si( classes of microstructure obtained by modification in the case of
hypoeutectic alloys. #lass 5 represents the super modified structure which is not repeatable.
MaJority of the modified castings will have structure of 0!4 type. %Modification ratings& is
generally used to assess the efficiency of the modifier. It is calculated by measuring the
percentage of each class of modification on a well!polished sample by optical metallographic
techniques. The variables that determine the microstructure are the type of modifier and its
amount impurities present in the metal free.ing rate and *i content in the alloy. =oth -a and
*r will produce all the si( classes of microstructure. There are some elements li"e > when
present in the metal will ma"e the modification difficult. In other words impurities present in
the metal will also play a role in getting modified structure. *imilarly the cooling rate or rate
of solidification of the melt also influences the modification. The modification process will be
assisted by higher cooling rate and hence less quantity of modifier will be sufficient in such
cases. It is natural that higher *i containing alloy need higher amount of modifier to get
particular class of modification. $n e(cess level of modifier will produce what is called an
over modified structure which will reduce the mechanical properties.
*r when used as a modifier may lead to the formation of $l
3
*r*i
1
intermetallic that
has deleterious effects on the mechanical properties. If the modified melt is "ept for a long
time the fading effect will become effective in which the levels of -a or *r in the melt will
come down with time. 2ading in case of -a is much faster than that in case of *r. The high!
resolution microscopic observation of high degree of twins in *i fibres of *r modified $l!*i
alloy threw light on the role of twins in the formation of fibrous eutectic *i. Impurity induced
twinning theory was also used to e(plain the mechanism of eutectic *i modification by a
chemical modifier and this became more widely accepted eutectic modification theory.
$ccording to this theory chemical modifiers are impurities that poison the growing atomic *i
layers by getting adsorbed on to surface steps producing "in"s and thus preventing further
attachments of *i atoms to the crystal. The adsorbed impurity atoms induce twinning by
altering the stac"ing sequence of atomic layers. The condition of growth twin was found to be
that the ratio of the impurity atom radius to the matri( atom radius should be 0.53.
The microstructure of hypereutectic $l!*i alloys consists of primary *i phase in a
binary eutectic. The presence of modifier such as *r ';.0<) shifts the eutectic composition
from 01.5< to around 03< as e(plained earlier and as a result the microstructure will contain
primary α dendrites in addition to primary *i and the eutectic. The presence of *r gives
complete eutectic structure also if the composition is Just above the eutectic composition.
+imited literature is available on the modification of both the eutectic and the primary
*i in hypereutectic alloys. The effect of *r addition on a > pre refined @:; alloy revealed
sufficient modification of eutectic at higher levels of *r ';.;@<). However the primary *i
particles were found to become coarser due to the addition of *r. Mn the other hand if -a is
used as a eutectic modifier instead of *r the primary *i is found to decrease in si.e and
become spheroidal along with well!modified eutectic *i. 7ery high concentration of -a 'of
the order of ;.0<) is required to achieve this structure. The combined addition of -a and *r
has found to result in a greater number of fine spheroidal *i particles than is obtained by
single -a treatment.