metallurgy of casting

8/12/2019 Metallurgy of Casting

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METALLURGY OF CASTING

Prof. B. S. MurtyDepartment 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 tinysolid until no liquid remains

Both conditions are met when the free energy of the particularphase 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.


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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 *+is the solid,liquid interfacial energy.

F%. (-ucleation barrier during solidification


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*olving the above equation for the overall free energy minimiation 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,Tand G/

0,T1. 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 stabilied '2ig. 5). The formation of dendritic

microstructure has to be suppressed particularly in directional solidification and one


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refining where compositional segregation is intended to be minimied. $ 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


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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.Sanchezand #-&- 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 freeing 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

freeing 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 aSolid Solution (lloy

Cooling Curve

F%. 0#ooling curve for the solidification of

single phase alloy


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F%. &1Effect of freeing 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


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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 sie. 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 sie reaches a minimum which is commonly referred

to as ultimate grain size). The time required to reach ultimate grain sie 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 sie 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 sie 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


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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 B1Ti25and B=23with 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 studieshave shown that the overall reaction could be e(pressed as follows.

01$l F @GB1Ti25 F 1GB=23 1GTi$l@ F GTi=1 F 4GB$l23 F B@$l25

$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 B1Ti25and B=23. However this practice was

later abandoned due to the difficulties in controlling the Ti and = contents high dross andlarge 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


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$ 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 beenused 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 sie and sie 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. $ higherreaction temperature '0;;;;#) during master alloy preparation is "nown to result in plate

li"e Ti$l@ particles while at lower temperatures '9;;;#) 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 coolingrate as in the case of chill casting the fla"y nature of the eutectic *i changes to fine fibrous


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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 freeing 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


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sensitivity loss by vaporiation 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 ! $l

percentage as well as its morphology in the case of eutectic $l!*i alloy. It has been observedthat 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 !phase by around

1@5< compared to the unmodified alloy. The increase in the amount of !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.Ithas been found that *r addition increases the volume percentage of the porosity pore sie 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 hasbeen 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 sie as well as shape is obtained. The under cooling produced by

the addition of small amounts of *b to eutectic $l!*i alloy showed the presence of !$l in themicrostructure. 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 sie

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


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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 ofhypoeutectic 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 freeing 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 suchcases. 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*i1 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

metallurgy of casting

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