Experiments were conducted on metal droplet undercooling, using Sn-25wt%Pb and Ni-34wt%Sn alloys. To achieve the high degree of undercooling, emulsification treatments were employed. Results show the fraction of supersaturated primary phase is a function of the amount of undercooling, as is the fineness of the structures. The solidification behavior of the tin-lead droplets during recalescence was analyzed using three different hypotheses; (1) solid forming throughout recalescence is of the maximum thermodynamically stable composition; (2) partitionless solidification below the T sub o temperature, and solid forming thereafter is of the maximum thermodynamically stable composition; and (3) partitionless solidification below the T sub o temperature with solid forming thereafter that is of the maximum thermodynamically metastable composition that is possible. The T sub o temperature is calculated from the equal molar free energies of the liquid solid using the regular solution approximation

Chu, M. G.

Flemings, M. C.

Macisaac, D. G.

Shiohara, Y.

English

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SOLIDIFICATION MECHANISM OF HICKLY UNDERCOOLED METAL ALLOYS
YUH SHIOHARA, MEN G. CHU, DAVID G. 1 4ACISAAC AND MERTON C. FLEMINGS
H.I.T., 77 Massachusetts Avenue, Cambridge, Hassachusetts, LISA -
ABSTRACT
This paper describes the results of experiments on
metal droplet undercooling, and presents a preliminary
analysis of the structures obtained. Work was conducted
with Sn-25wt%Pb and Ni-34wt%Sn alloys. To achieve the re-
quired high degrees of undercooling, emuls i fication treat-
ments were utilized following techniques described-by
Perepezko and co-workers.
Experiments show the fraction of supersaturated
primary phase is found to be a function of the amount of
undercooling, as is the fineness of the structures. The
solidification behavior of the tin- lead droplets during
recalescence was analyzed using three different hypotheses;
l) solid forming throughout recalescence is of the maxfWum
thermodynamically stable composition, 7) partitionless
solidification below the "To" temperature, and solid form-
ing thereafter is of the maximum thermodynamically stable
composition, and 3) partitionless solidification below the
"To" temperature with solid forming thereafter that is of
the maximum thermodynamically metastable composition that
is possible.
INTRODUCTION
Undercooling of liquid metal is a common phenomenon, but the amount of
undercooling is usually restricted by the presence of heterogeneous nucleation
catalysts. The homogeneous and heterogeneous nucleation theory was first
applied to metallic systems by Turnbull and his co-workers [1]. The critical
(maximum) undercoolings obtained were approximately equivalent to 0.18 T m , where
Tm is the melting point of the metal in the absolute temperature scale.
Recently, the droplet technique has been modified and improved by
Perepezko, Loper and co-workers [2,3]. Finel y dispersed droplets of low melt-
ing temperature alloys, emulsified in an organic oil, have shown undercoolings
in excess of 0.3 Tm prior to nucleation.
In this work, high undercoolings were generated by the emulsification
method and resulting microsegregation and the microstructure studied in Sn-
25wt2Pb alloy . The technique was also extended for the first time to u high
melting point nfekel base alloy.
EXPERIMENTAL PROCEDURE
The tin-lead alloy was prepared in a vacuum melter from high purity
(99.99997) metals. Droplet dispersions of the alloy were produced in a resis-
tance furnace by emulsifying a mixture of liquid metal and organic carrier
fluid (Polyphenyl-ether), contained a Pyrex crucible, with a high speed
(3800 rpm) shearing device. The atmosphere in the furnace was purged with pure
argon gas. Residual amounts of oxygen in the presence of isophthalic acid
served as an oxidant to produce a chin oxide surface film, which prevents the
`..- 4--4
it ff. °19
e
.OF	 LMQUA
droplets from agglomerating. The temperature was kept at 25°C above the melt-
ing, point during sharing, which continued for 40 minutes, producing fine drop-
lets dispersed in the oil. The emulsified samples were transferred to sealed
aluminum pans for analysis in a differential thermal analysis (DTA) or a differ-
ential sewing colorimeter (DSC) apparatus. The emulsified samples were re-
moved from the aluminum pans and ultrasonically cleaned in preparation for obser-
vation of cross sectional microstructures by scanning electron microscope (Slit).
The emulsification method was modified for use with a nickel-tin allov. Dis-
persions of the alloy in an oxide glass was carried out in an alumina crucible,
under an inert atmosphere, with a mixing device also made of alumina. Cinder-
codlings of large single droplets and droplet dispersions were measured by DTA.
RESULTS
	
T 	 T  TL
Tin-Lead A112y. The emulsified
	 i	 1
droplets, observed with SEM, were 	 10°C/min=
distributed in the size range of 10-20 	 Heating
Um. A representative DTA result of
about 15 mg of.alloy and 20 mg of oil
is presented in Fig. 1. The thermal
	 30 mca
heating; curve of Sn-25wt2Pb ( = 16at%Pb)	 i
allay shows two endothermic peaks due
to melting. The first peak indicates
the melt q, of eutectic regions. The
	
Cooling
second peak corresponds to the melting	 •	 1	 A t
of primary tin rich phase. During	 80 100 120 140 160 180 200 220 240
cooling, one sharp exothermic peak due
	
Temperature(°C)to solidification was observed. The
amount of undercooling of most drop-
lets is defined by the onset of this
sharp peak. The shaded area in the 	 Fig. 1. Representative DTA results
figure is related to the latent heat of a Sn-25wt%Pb droplet emulsion.
of fusion for an alloy.
Figure 2 shows the cross sectional microstructure of a typical droplet
which was solidified at a sample cooling; rate of 10°C/min. Undercooling was
measured as 80°C. The microstructure observed consists of a tin rich region
surrounding a smaller lead rich region. The volume fraction of the lead rich
region was estimated at 77. This fraction was found to decrease with increas-
ing; cooling rate.
Follo.4ing
 the same heating and cooling treatment, a specimen was re-
heated to 1°C higher than the eutectic temperature, held for 10 minutes (pri-
mary
	remained as a solid), and then cooled at
- the 10°C/min cooling rate.
The undercooling for the eutectic solidification was measured as 6-9°C by DTA.
The microstructure of this sample is shown in Fig. 3. Note that the lead rich
eutectic now surrounds the primary tin rich phase.
Nickel-Tin A l l oy. The sizes of the dispersed droplets produced by Ni-
14w Sn ran '
	
Ped from 5-3000 um. Fig. 4 show SErt micrographs of several of
these droplets. The predominant phase in all these structures was analyzed as
the hypo-eutectic (tin rich) phase. Figs, 4F, and 4b show strictures of coarse
droplets of the same size (2 mm), nucleated at different temg• eratures. The
finest structure corresponds to the highest undercooling. Vie droplets in
Figs. 4c and 4d are much smaller in size (700 Pm and 60 um d;.ameter respective-
ly); both are from a single dispersed droplet DTA run in which the undercooling
was measured as a single peak. Note the structure is finer in the smaller of
these droplets and both these structures are much finer than those in Figs. 4a
and 4b.
(c) (d)
(b)
ORIGINAL Fnwy 1';o- 	 3 L
OF PCOR QUAL"
Fig. 2. Cross sectional microstruc-
ture of a Sn-25wt%Pb droplet after
solidification at a cooling rate of
10 0 C/min and with 80°C of undercooling.
Fig. 3. Cross sectional microstruc-
ture of a Sn-25wt% Pb droplet in which
the eutectic region was remelted at
1°C higher than the eutectic tempera-
--:re and then cooled at 10°C/min.
Fig. 4. Cross sectional microstructure of Ni-34wt%Sn droplets.
(a) Droplet diameter (D) = 2 mm, undercooling (AY) = 88°C.
(b) D = 2 mm, AT = 177°C.
(c) D = 0.7 mm, AT = 219°C.
(d) D = 0.06 mm, AT = 219°C.
4`
OF POM QUM*
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a H
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ai
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	 CS1C.K ',$ crib C.,	 CsR Cs!! Co CM CE
COMPOSITION -
	
COMSITION
	
T COMPOSITION
Fig. 5. Schematic illustration of the solidification paths for three solidi-
fication models.
DISCUSSION
In order to interpret the structures and segregation obtained in under-
cooled specimens we have carried out detailed calculations of solidification t
behavior using the three models to be discussed bel^-)w. Complete calculations
and extensions of these models are described in a separate paper (4). Assump-
tions common to all three models are 1) no solid diffusion, 2) no remelting
during recalescence, 3) adiabatic conditions during recalescence and 4) negli-
gible thermal gradients in the drop-
lets during solute redistribution,
and after recalescence, solidifiea-
tion proceeds according to the
Scheil equation.
In Case 1, solid forming throughout
solidification is assumed to be of the 15
maximum thermodynamically stable compo- 4
.,
sition.	 Hence initial solid forms of
compositions along the solidus from Csn to
at the nucleation temperature to CsR at S	 10
the maximum recalescence temperature.
o
If diffusion in the liqu id is essen-
tially complete, liquid composition 8	 5
follows the line shown in Fig. 5 for Co`i
to CLR during recalescence.	 After re-
o
calescence, solidification o ccurs by
0the Scheil a nation with the	 rimaryq	
pU	 0 0.2 0.4 0.6 0.3 1.0solid composition reaching a maximum
of CsN and the liquid reaching a maxi- 	 Fraction solid, fs
mum of eutectic composition, C F (5].
Fig. 6, curve 1, shows the solute
redistribution expected across	 Fig. 6. Calculated solute redis-
a cell or dendrite of Sn-16at2Pb alloy 	 tribution profiles for three
undercooled 80°C and solidified accord- 	 solidification models, Sn-16at%Pb.
ing to these assumptions.
In Case 2, solidification is assumed partitionless between TN and To so
solid composition is Co. The "To" temperature is calculated from the equal
molar free energies of the liquid and the solid using the regular solution
approximation. The dependence of the temperature on the molar free energy is
S
`^R1(1NRI. PACE 1
OF KYUR UAL11Y	 5
calculated from the Gibbs-Helmholtz relationship. Above To, solidification
'	 is as in Case 1.
In Case 3, solidification below To is partitionless as in Case 2. Above TO*
the solid composition at the interface is the maximum thermodynamically possi-
ble solubility. This composition, Cs, is calculated from the intersection of
the molar free energy curve for the solid and tangent line to the molar free
energy curve for the liquid at the composition CL, as described by Cahn [6].
On the basis of microstructures obtained to date, it is not possible to
distinguish between the three mechanisms discussed above. Composition calcu-
lations based on lineal analysis of volumes of the two phase present in the
photomicrograph of Fig. 2 indicate the composition of the large tin rich re-
gion is just below Co (about 22wt%Pb) and the small lead rich region is about
65wt%Pb. If Case 1 applles, liquid must have been-entrapped during growth of
the large tin rich region. If the small region represents solidification of
eutectic liquid after its recalescence, its amount (estimated volume fraction
.07) is much lower than that of any of the three cases considered, in which
predicted fraction eutectic varied from about 0 . 45 to 0.6, and its composition
is much higher than would be expected. It is clear that other factors must be
considered in predicting solute redistribution in these samples, including
perhaps heat flow 'during and after recalescence, remelting and ripening, eutec-
tic undercooling, and metal shrinkage or expansion during solidification.
Fig. 3, obtained by holding 10 minutes at l *C above the eutectic tempera-
tore and then cooling shows two significant effects. One of these is the
dramatic change in shape of the primary solid by ripening and the other is the'
coarse lamellar type eutectic structure typical of eutectic solidification at
low undercoolings [7]. Both those effects may need to be considered in a full
analysis of the original undercooled structure.
The fineness of the Ni-Sn structure, as well as the "eutectic" structure
of the Sn-Pb alloy is indicative of extremely rapid solidification rates. The
structure of the large Ni-Sn droplet, solidified at high undercooling, is only
about 5Wn while that in the 601im diameter particle is less than about 1pm even
though the cooling rate of the specimens was only about 50°C/min. Local heat
extraction from the individual droplets into the surrounding fluid must be an
important factor influencing these structures.
ACKNOWLEDGEMENT
This work was supported by NASA Grant No. NSG-7645.
REFERENCES
1. D. Turnbull and R.E. Cech, J. Applied Physics, 21, 805 (1950).
2. J.H. Perepezko, D.H. Rasmussen, I.E. Anderson and C.R. Loper, Jr.,
Proceedings of International Conference on Solidification, The Metal
Society, 1!A (1977).
3. J.H. Perepezko and I . E. Anderson, TM$-AIME Symposium on Synthesis and
Properties of Metastable Phases, T.J. Rowland and E.S. Machlin eds.
In press.
4. M.C. Flemings, Y. Shiohara, M.G. Chu and D.G. ?iaclsaac, to be published.
5. T.Z. Kattamis, Zeitschrift fur Metallkunde, 61, 856 (1970).
6. J.C. Baker and J . W. Cahn, Solidification (American Society for Metals,
Metals Park, Ohio, 1971) p. 23.
7. T.Z. Kattamis and H.C. Flemings, Trans. AIHE, 236, 1523 ( 1966).


Solidification mechanism of highly undercooled metal alloys

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