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Feeding to castings from a riser in a pasty state
Materials Science and Engineering, 96 (1987) 259-265
259
Feeding to Castings from a Riser in a Pasty State C. Y. LIU, K. MURAKAMI and T. OKAMOTO
The Institute of Scientific and Industrial Research, Osaka University, 8-1 Mihogaoka, Ibaraki, Osaka 567 (Japan) S. GODA
Faculty of Science and Technology, Kinki University, 3-4-1, Kowakae, Higashiosaka, Osaka 577 (Japan) (Received June 10, 1987)
ABSTRACT
The a m o u n t o f porosity and the degree o f macrosegregation are measured in A l 4.5wt. %Cu alloy ingots solidified unidirectionally upwards from the chill face. When low melting p o i n t liquid metal is added onto the top o f solidifying ingots just before the top region becomes in a pasty state, the a m o u n t o f porosity is smaller and the degree o f macrosegregation larger than those in an ingot to which no liquid metal has been added. This is because the addition does n o t increase the resistance to liquid f l o w which, in the case o f no addition, would come from a capillary force resulting from the surface tension at the top free surface o f the interdendritic liquid and prevent the m o v e m e n t o f the interdendritic liquid. The a m o u n t o f porosity and the degree o f macrosegregation in the ingots with and w i t h o u t the addition were also c o m p u t e r simulated and are compared with the experimental results.
ing only from viscous friction. The effect of capillary pressure shown in Fig. l(b) on the interdendritic liquid flow will be important, especially in the solidification of an alloy with a wide freezing range. In this case, when the riser solidifies to become a pasty zone, the solidification of the casting may n o t be totally complete; then further solidification shrinkage of the casting cannot be easily compensated for by feeding liquid from the riser. The aims o f the present work are to elucidate the effect of the capillary pressure on the formation of porosity and macrosegregation and to estimate them by means of computer simulation of the solidification. 2. EXPERIMENTAL PROCEDURE A cylindrical mould made of a heat-insulating brick was used to make A1-4.5wt.%Cu
freesurfQce liquid of
ol[oy
~
1. INTRODUCTION Three of the present authors have shown [1, 2] that, when the top free surface of liquid flowing downwards through a porous medium came into the medium, the liquid flow w a s made difficult because of a capillary pressure arising from the surface tension of the liquid. This p h e n o m e n o n will also take place in casting practice when the top of a riser becomes in a pasty state, as shown schematically in Fig. l(b). Figure l(a) shows a hypothetical case in which there is no capillary pressure acting on the liquid and so the liquid in the interstices of solid network flows down relatively easily with the resistance to flow aris0025-5416/87/$3.50
solid-
flOW
(o)
flow rate = 0
(b)
Fig. 1. Schematic drawing of the free surface of the interdendritic liquid in the two models: (a) model I in which no capillary pressure is assumed to be present at the liquid free surface and the liquid can move relatively easily through the interdendritic space; (b) model II in which capillary pressure appears at the liquid free surface and prevents interdendritic liquid flow.
© Elsevier Sequoia/Printed in The Netherlands
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perature of the AI-4.5wt.%Cu alloy was added onto the ingot top. For ingot C, pure lead melt was added instead of the eutectic melt. The solidified ingot was sectioned longitudinaUy and one haft of it was polished and etched for observation of the macrostructure. From the other haft of the ingot, specimens for examination of the macrosegregation and the a m o u n t of porosity were taken at various distances from the b o t t o m of the ingot. The a m o u n t of porosity in the specimens was determined from the density of the specimens and the density of hot isostaticaUy pressed specimens which were free from porosity and contained 3.1-6.1wt.% Cu. Hot isostatic pressing was done under a pressure of 51 MPa at 833 K.
3. COMPUTER SIMULATION
alloy melt solidify unidirectionally upwards as shown in Fig. 2. The height, the inside diameter and the outside diameter of the mould were 140 mm, 48 mm and 120 m m respectively. An aluminium chill plate 18 mm thick was set at the b o t t o m of the mould. The mould was dried at 423 K for 10 h. A14.5wt.%Cu alloy melt was made by melting 99.9 wt.% Al and A1-51.8wt.%Cu master alloy in a graphite crucible. The melt was degassed and then poured into the mould with a superheat of 130 K. To ensure that the solidification of the melt proceeds unidirectionally upwards from the chill face to the top, a graphite plate 3 mm thick was put on the top of the mould and then was heated using a gas burner. At the same time the a h m i n i u m chill was cooled by blowing compressed air on it. These operations were carried on during the whole period of the solidification. The temperatures in the ingot were measured with five c h r o m e l - a h m e l thermocouples 0.3 mm in diameter located on the longitudinal axis of the ingot at 5, 30, 55, 80 and 105 mm from the bottom. The temperatures of the chill and the mould, and the temperature in the space between the graphite plate and the ingot top were also measured. Three ingots A, B and C were cast. Ingot A was solidified in the manner mentioned above. Just before the top of ingot B reached a pasty state, A1-CuA12 eutectic alloy melt held at a temperature of 10 K below the liquidus tem-
Macrosegregation and the a m o u n t of porosity were numerically calculated for the ingots obtained in the present work, which consisted of columnar dendrites growing upwards from the bottom. The method of calculation was the same as that described by Kubota et al. [ 3 ] except that the formation of porosity was considered in the present work. For the purpose, an additional equation to describe the formation of porosity was used, which held in the pasty zone:
V~AFv = -- aV~ A fL -- ~, fLVSj At
(1)
J
The flow rate U in the pasty zone can therefore be obtained from the following equation.
Z
i PLfLUSj
AtS1 ~
Ps(h + ~m) CLi(ko--1) (Phi-- Ps)Pq~
-cLi)
(h + pu~
CpL(P~ -- ps)V~(T-- T~)
(Pu-- Ps)(PsfslCps + P~fuCpL) Vim( CL--Cu) ~,
PufLi _ (p=psV a f v
8j(Ti -- T~) At Z J l/a,~- Axdk~- Axdkj (2)
261
PsV~{ =
~'=1+
-
(Psf~Cps+~pLifuCpL)mc (ko-1)j
(h + ~m)psC~(ko-
1)
(4)
g N = 8 . 8 X 10-19dll'3d22"4fL3"2
4. RESULTS
PL~(PI~ - - P s )
Calculation was carried o u t on the basis of the following two models. In the first model (model I) which corresponded to ingots B and C, the capillary pressure was assumed to be negligible. According to this model, the interdendritic liquid at the t o p of the original b o d y of the ingot flows down until solidification of the region is completed. In the second model {model II) which corresponded to ingot A, it was assumed that the interdendritic liquid stopped flowing because of the capillary pressure after the fraction solid at the ingot t o p exceeded 0.2. This assumption resulted from the observation by O k a m o t o [4] that secondary arms growing from neighbouring columnar dendrites began to t o u c h each other at a solid fraction between 0.16 and 0.22. The following initial and b o u n d a r y conditions were adopted. At time zero in the bulk liquid, the temperature of the melt is uniform and equal to the pouring temperature and the melt is quiescent. The initial temperature of the mould is 353 K and the ambient temperature of the mould is 298 K. The flow rate is zero on the mould wall. The temperature in the space b e t w e e n the graphite plate and the ingot top is 1123 K. The physical and thermal properties of A14.5wt.%Cu alloy, mould and aluminium chill are given in Table 1 and in Appendix A. The permeabilities Kp and K s for liquid flow perpendicular to the columnar dendrites and normal to the columnar dendrites respectively are as follows [ 5 ] : K p : 6.2 X lO-3d12"2d2-1"7fL3"2
(3)
The measured cooling curves in ingot A at three locations 5, 55 and 105 mm from the chill face are shown by the full lines in Fig. 3. The d o t t e d lines in the figure are the cooling curves calculated on the basis of model II and are in good agreement with the experimental lines. Figure 4 shows the measured cooling curves and the calculated curves for ingot B. It is apparent from Fig. 3 or Fig. 4 that there is a positive temperature gradient from the b o t t o m to the t o p of the ingots during the whole period of solidification. The arrow in Fig. 4 shows the time at which A1-CuA12 eutectic melt was added onto the ingot top. Only for a short period after the addition did the temperature slightly decrease at the location of thermocouple 3 and in the upper portions. The macrostructures in the longitu-
1073 /
i
973
,,=,
873
~
........
•
5 .....
........ 1
773
673
, 300
i 600
, 900
, 1200
, 1500
1800
TIME ( s )
Fig. 3. Measured ( ) and calculated (..... ) cooling curves in ingot A. The numerals 1, 3 and 5 represent the thermocouples in Fig. 2. T e is the eutectic temperature of the AI-4.5wt.%Cu alloy.
TABLE 1 Values of the density, thermal conductivity and specific heat used in the calculation
AI-4.5wt.%Cu, liquid AI-4.5wt.%Cu, solid AI-4.5wt.%Cu, eutectic solid Brick Aluminium chill
p
~
ce
(kg m -3)
(W m -1 K -1)
(kJ kg -1 K -1)
PL = PO + h(Cb -- Co) + ~(T-- TO)
138.1 138.1 138.1 138.1 138.1
1.088 0.896 0.896 0.837 1.075
2620 3400 670 2700
262
dinal sections of ingots A, B and C are shown in Fig. 5. The densities of the porosity free hot isostatically pressed specimens are plotted against the copper content in Fig. 6. This relation is formulated as follows: p = 2700 + 22 × wt.% Cu
(5)
where p is the density. The amount of porosity in the ingots was obtained from the density measurement, as is shown in Fig. 7. The data for ingots B and C are neglected near the ingot tops because
the added low melting point liquid metal penetrated into the ingot over some distance from the tops. It can be seen that the addition of the low melting point liquid metal onto the ingot top resulted in a significant decrement in the amount of porosity compared with the case of no addition. The amount of porosity calculated on the basis of model II, which corresponds to ingot A, becomes larger with increasing distance from the chill surface as shown by the broken line in Fig. 7. The amount of porosity calculated from model I corresponding to ingots B and C is essentially zero throughout the ingots. The calculated
1073 3.0 xl03
A
973 2.9
I 873
/I
28
,,=,
Te = 821 K
t-
~
o ~~
I +o ~ + + ~ +
,
o++
++
773 2.7, 673 0
i 300
tPl, +
900
600 TIME
L
i
1200
1500
1800
( s )
Fig. 4. Measured ( ) and calculated ( . . . . . ) cooling curves in ingot B. tp is the time at which AI-Cu eutectic ahoy liquid was added onto the ingot top.
2,6 0 0
' 1.0
' 2.0
' 3.0 Cu
' 4.0
' 5D
CONTENT (
wt
' 6.0
7.L0
' &O
' 9.0
150
pct)
Fig. 6. Densities of AI-Cu alloys free from porosity: o, measured; e, ref. 6.
mm
100
50
Fig. 5. Macrostructures on the longitudinal sections of unidirectionally solidified AI-4.5wt.%Cu alloy ingots: (a) ingot A; (b) ingot B; (c) ingot C.
263
4.0
3.0
e°
3 F--
o no Q.
2.0
::::t 14.0[-
I---
zt ~
6.0
b--
z
ot 9
1.0
5.0 ~-%~- ~ ~,_ _---_--__ "__ _ ~ * _ ~ _
...........
3
c)
0.0
O
I
I
I
I
I
20
40
60
80
100
120
4.0
3.0
DISTANCE FROM CHILL ( m m )
0
2'0
40
I 60
, 100
J 80
120
DISTANCE FROM C H I L L ( m m )
Fig. 7. Amount of porosity in ingots A (o), B (e) and C (c~): - - -, value calculated from model II.
Fig. 9. Macrosegregation in ingot B: o, experimental results;--, results calculated f r o m m o d e l I; - - -, results calculated f r o m m o d e l II.
5.5 o
~" ,...;
18.0 /
5.0
17.0 II-Z
,,,
/-,.5
I,Z
o (.9
v FZ
4.0 3.5
LU
13.OJ~
ou
T 60 r
)-Z
0
210
41O
, 60
I 80
= 100
120
DISTANCE FROM CHILL ( m m )
Fig. 8. Macrosegregation in ingot A: e, experimental r e s u l t s ; - - , results calculated from model I; - - -, results calculated from model II.
a m o u n t o f p o r o s i t y in t h e ingots is smaller t h a n t h e e x p e r i m e n t a l value. T h e d i f f e r e n c e s between the calculated and the measured values f o r t h e t h r e e ingots agree w i t h e a c h other. T h e m e a s u r e d m a c r o s e g r e g a t i o n in i n g o t A is s h o w n b y closed circles in Fig. 8. T h e c o p p e r c o n c e n t r a t i o n d e c r e a s e s w i t h increasing dist a n c e f r o m t h e chill s u r f a c e , i.e. inverse segreg a t i o n occurs. Also in t h e figure are s h o w n t h e c a l c u l a t e d results f r o m m o d e l s I a n d II. T h e m a c r o s e g r e g a t i o n is larger in m o d e l I t h a n in m o d e l II. T h e d e g r e e o f m a c r o s e g r e g a t i o n c a l c u l a t e d f r o m m o d e l I I agrees q u i t e well w i t h t h e e x p e r i m e n t a l values. T h e m a c r o s e g r e g a t i o n in i n g o t B is s h o w n in Fig. 9. I n this case, g o o d a g r e e m e n t is o b t a i n e d b e t w e e n e x p e r i m e n t a n d c a l c u l a t i o n o n t h e basis o f m o d e l I. I n Fig. 10 t h e c o p p e r a n d lead c o n -
14.0~ °
•
•
°
s.o F ~ - ~ . . _ . , _ ........ - - ~ _ : : : : : : U
..........
40k
:o/0
,
,
,
20
40
60
80
100
120
DISTANCE FROM CHILL ( r a m )
Fig. 10. Distribution of copper and lead in ingot C: o, experimental results for copper; o, experimental results for lead; - - . , results calculated from model I; - - -, results calculated from model II.
t e n t s in i n g o t C are p l o t t e d . I t can be seen t h a t lead p e n e t r a t e d i n t o t h e ingot. 5. DISCUSSION T h e solidification o f all t h e ingots in t h e present work proceeded upwards from the chill f a c e . T h e liquid f l o w t o c o m p e n s a t e f o r solidification s h r i n k a g e in t h e ingot, t h e r e f o r e , s h o u l d o c c u r vertically d o w n w a r d s f r o m t h e i n g o t t o p . I n i n g o t A, t o w h i c h n o l o w m e l t i n g p o i n t liquid m e t a l w a s a d d e d , t h e f e e d i n g process c a n be divided i n t o t w o stages. In t h e first stage, t h e r e is a b u l k liquid r e g i o n a b o v e t h e
264
pasty zone; the bulk liquid is supplied into the pasty region to feed solidification shrinkage in this region. In this case the liquid flows easily against liquid flow resistance coming only from viscous friction in the pasty zone. In the second stage where the ingot top becomes in a pasty state, the capillary pressure appears because of the surface tension of the interdendritic liquid and acts as an additional resistance to the interdendritic liquid flow, accelerating the formation of shrinkage porosity and blow holes. This situation is also explained as follows. The pressure of the interdendritic liquid just beneath the free surface is smaller (by the capillary pressure) than the atmospheric pressure. The pressure of the liquid in the ingot is therefore much smaller than that in the case where there is no capillary pressure. As a result, pores are easily formed in the ingot. The t o p surface of ingot A was rather smooth after solidification had been completed, and there was no dendrite arm protruding from the top surface to the air. This fact indicates that the capillary pressure is very large and that no interdendritic liquid could flow down in the second stage in ingot A. The results in Fig. 7 show that the a m o u n t of porosity decreases largely by the addition of low melting point liquid alloy onto the ingot top. The liquidus temperature of the top layer of ingot B was lower than that of A1-4.5wt.%Cu alloy because of the addition of eutectic alloy melt. The capillary pressure, therefore, began to arise in a much later stage of the solidification than in the case of ingot A. At that time, the fraction of liquid in the original part of the ingot was small and the amount of porosity should be small. The lead layer at the top of ingot C, shown in Fig. 5(c), was liquid even when the temperature at the ingot top fell below the eutectic temperature of the A1-CuA12 eutectic alloy system. In this case, the capillary pressure arises because of the interfacial tension b e t w e e n the interdendritic alloy liquid and liquid lead, which is much smaller than that from the surface tension. Feeding liquid from the ingot top is therefore much easier in ingots B and C than in ingot A. This suggests that a large portion of the porosity in ingot A was generated after the t o p of the ingots solidified to become a pasty zone. The fact that the amount of porosity calculated as the basis of model I
or II for each ingot is smaller than the experimental value may be ascribed to the evolution of hydrogen gas during solidification of the ingots. Inverse segregation arises because the soluterich interdendritic liquid flows towards the chill surface to compensate for the solidification shrinkage [7, 8]. The capillary pressure appearing at the t o p of ingot A makes the flow of solute~enriched interdendritic liquid more difficult and decreases the macrosegregation. The resultant segregation in ingot A agrees quite well with the calculated value in which the liquid flow rate is assumed to be zero after the ingot t o p had become in a pasty state. In contrast, when low melting point liquid alloy was added onto the t o p of ingot B, feeding to the ingot was made relatively easy. The segregation can be successfully predicted on the assumption of complete flow-back of soluteenriched interdendritic liquid as shown in Figs. 9 and 10. It can be said from the results in Figs. 8 - 1 0 that, the larger the amount of porosity, the smaller the degree of segregation is.
An oxide film is easily formed on the surface of liquid aluminium in air. In fact, the top surface of ingot A was covered with an oxide film. Kahl and F r o m m [9] showed that the strength of the oxide film on the liquid surface of pure aluminium and aluminium alloys is much larger than the surface tension of the corresponding liquid. In such a case, after the ingot top becomes in a pasty state, the interdendritic liquid flow would be much more difficult than in the case where the liquid-free surface is free from an oxide film. It is concluded that, when the free surface of a riser or risers becomes in a pasty state, since the capillary pressure or the surface oxide film makes the feeding of liquid to ingots and castings difficult, it accelerates the formation of porosity and also reduces the macrosegregation, in the case of alloys with a wide freezing range.
6. CONCLUSIONS
A1-4.5wt.%Cu alloy ingots were solidified unidirectionally upwards and, just before the ingot t o p became in a pasty state, A1-CuA12 eutectic alloy melt or lead melt was added onto the t o p to keep the t o p region in a bulk
265 liquid state during solidification of the original part of the ingot. When the ingot top becomes in a pasty state, the free surface of the liquid descends into the interdendritic space and causes a capillary pressure to develop which increases the resistance to interdendritic liquid flow. The addition of the low melting point liquid keeps the resistance to interdendritic liquid flow at the same level as that from only viscous friction, because it prevents the occurrence of capillary pressure in the liquid. The interdendritic liquid can therefore easily move downwards into the pasty region to compensate for the solidification shrinkage. As a result, the degree of macrosegregation becomes larger and the a m o u n t of porosity becomes smaller in these ingots than in an ingot to which no liquid metal was added. The macrosegregation in a melt~added ingot can be successfully predicted by computer simulation on the assumption that there is no resistance to the interdendritic liquid flow at the ingot top, whereas t h a t in a no-melt-added ingot can be predicted by assuming t h a t the interdendritic liquid stops flowing after the fraction solid at the ingot top exceeds 0.2. The difference in the a m o u n t of porosity between the ingots with and w i t h o u t the melt addition is qualitatively predicted.
4 T. Okamoto, Tetsu To Hagan$, 58 (1982) 1302. 5 K. Murakami, S. Shiraishi and T. Okamoto, Acta Metall., 32 (1984) 1423. 6 E. A. Brandes, Smithell's Metals Reference Book, 6th edn., Butterworths, London, 1983, p. 14-14. 7 J. S. Kirkaldy and W. V. Youdelis, Trans. MetaU. Soc. AIME, 212 (1958) 833. 8 M. C. Flemings and G. E. Nereo, Trans. Metall. Soc. AIME, 239 (1967) 1448. 9 W, Kahl and E. Fromm, Z. Metallkd., 75 (1984) 957.
APPENDIX A: NOMENCLATURE CT.
solute concentration of liquid
Cp~. specific heat of liquid
This work was supported by the research project of Institute of Scientific and Industrial Research, Osaka, on the development of new materials for energy. The authors wish to t h a n k Professor M. Koizumi, Professor Y. Miyamoto and Dr. K. Suganuma for their help in hot isostatic pressing, and Mr. M. Ona for performing the experiments.
Cps specific heat of solid Co = 4.5 wt.%, solute concentration of liquid at the liquidus temperature dl primary arm spacing d2 -- 50 m, secondary arm spacing fL volume fraction of liquid fs volume fraction of solid Fv a m o u n t of porosity h = 26.3 m -s (wt.%) -1, constant in Table 1 h a heat transfer coefficient H -- 397.5 kJ kg-1, heat of fusion k 0 = 0.172, equilibrium distribution coefficient K N permeability normal to columnar dendrites Kp permeability parallel to columnar dendrites m = -- 3.39 K (wt.%y 1, liquidus slope S area of the side of an element t time T temperature To = 918 K, liquidus temperature Ax distance from a nodal point to the grid line U flow rate V volume of an element
REFERENCES
a
1 C. Y. Liu, K. Murakami and T. Okamoto, Acta Metall., 34 (1986) 159. 2 C. Y. Liu, K. Murakami and T. Okamoto, Acta Metall., 34 (1986) 1173. 3 K. Kubota, K. Murakami and T. Okamoto, Mater. Sci. Eng., 76 (1986) 67.
PL Ps P0
ACKNOWLEDGMENTS
Greek symbols
change in volume on solidification = -- 0.3 kg m -3 K -1, constant in Table 1 thermal conductivity density of liquid density of solid = 2440 kg m -a, density of liquid at the liquidus temperature
259
Feeding to Castings from a Riser in a Pasty State C. Y. LIU, K. MURAKAMI and T. OKAMOTO
The Institute of Scientific and Industrial Research, Osaka University, 8-1 Mihogaoka, Ibaraki, Osaka 567 (Japan) S. GODA
Faculty of Science and Technology, Kinki University, 3-4-1, Kowakae, Higashiosaka, Osaka 577 (Japan) (Received June 10, 1987)
ABSTRACT
The a m o u n t o f porosity and the degree o f macrosegregation are measured in A l 4.5wt. %Cu alloy ingots solidified unidirectionally upwards from the chill face. When low melting p o i n t liquid metal is added onto the top o f solidifying ingots just before the top region becomes in a pasty state, the a m o u n t o f porosity is smaller and the degree o f macrosegregation larger than those in an ingot to which no liquid metal has been added. This is because the addition does n o t increase the resistance to liquid f l o w which, in the case o f no addition, would come from a capillary force resulting from the surface tension at the top free surface o f the interdendritic liquid and prevent the m o v e m e n t o f the interdendritic liquid. The a m o u n t o f porosity and the degree o f macrosegregation in the ingots with and w i t h o u t the addition were also c o m p u t e r simulated and are compared with the experimental results.
ing only from viscous friction. The effect of capillary pressure shown in Fig. l(b) on the interdendritic liquid flow will be important, especially in the solidification of an alloy with a wide freezing range. In this case, when the riser solidifies to become a pasty zone, the solidification of the casting may n o t be totally complete; then further solidification shrinkage of the casting cannot be easily compensated for by feeding liquid from the riser. The aims o f the present work are to elucidate the effect of the capillary pressure on the formation of porosity and macrosegregation and to estimate them by means of computer simulation of the solidification. 2. EXPERIMENTAL PROCEDURE A cylindrical mould made of a heat-insulating brick was used to make A1-4.5wt.%Cu
freesurfQce liquid of
ol[oy
~
1. INTRODUCTION Three of the present authors have shown [1, 2] that, when the top free surface of liquid flowing downwards through a porous medium came into the medium, the liquid flow w a s made difficult because of a capillary pressure arising from the surface tension of the liquid. This p h e n o m e n o n will also take place in casting practice when the top of a riser becomes in a pasty state, as shown schematically in Fig. l(b). Figure l(a) shows a hypothetical case in which there is no capillary pressure acting on the liquid and so the liquid in the interstices of solid network flows down relatively easily with the resistance to flow aris0025-5416/87/$3.50
solid-
flOW
(o)
flow rate = 0
(b)
Fig. 1. Schematic drawing of the free surface of the interdendritic liquid in the two models: (a) model I in which no capillary pressure is assumed to be present at the liquid free surface and the liquid can move relatively easily through the interdendritic space; (b) model II in which capillary pressure appears at the liquid free surface and prevents interdendritic liquid flow.
© Elsevier Sequoia/Printed in The Netherlands
260 BURNER
GRAPHITE PLATE
THERMOCOUPLES
i ,
:
i
,.,
•
4~
:,v " ,
BRICK M O U L D - --;--- "." i
4
•
. ,//~://-/
ALUMINIUM CHILL
AIR Fig. 2. Cylindrical mould for unidirectional solidification of AI-4.5wt.%Cu alloy.
perature of the AI-4.5wt.%Cu alloy was added onto the ingot top. For ingot C, pure lead melt was added instead of the eutectic melt. The solidified ingot was sectioned longitudinaUy and one haft of it was polished and etched for observation of the macrostructure. From the other haft of the ingot, specimens for examination of the macrosegregation and the a m o u n t of porosity were taken at various distances from the b o t t o m of the ingot. The a m o u n t of porosity in the specimens was determined from the density of the specimens and the density of hot isostaticaUy pressed specimens which were free from porosity and contained 3.1-6.1wt.% Cu. Hot isostatic pressing was done under a pressure of 51 MPa at 833 K.
3. COMPUTER SIMULATION
alloy melt solidify unidirectionally upwards as shown in Fig. 2. The height, the inside diameter and the outside diameter of the mould were 140 mm, 48 mm and 120 m m respectively. An aluminium chill plate 18 mm thick was set at the b o t t o m of the mould. The mould was dried at 423 K for 10 h. A14.5wt.%Cu alloy melt was made by melting 99.9 wt.% Al and A1-51.8wt.%Cu master alloy in a graphite crucible. The melt was degassed and then poured into the mould with a superheat of 130 K. To ensure that the solidification of the melt proceeds unidirectionally upwards from the chill face to the top, a graphite plate 3 mm thick was put on the top of the mould and then was heated using a gas burner. At the same time the a h m i n i u m chill was cooled by blowing compressed air on it. These operations were carried on during the whole period of the solidification. The temperatures in the ingot were measured with five c h r o m e l - a h m e l thermocouples 0.3 mm in diameter located on the longitudinal axis of the ingot at 5, 30, 55, 80 and 105 mm from the bottom. The temperatures of the chill and the mould, and the temperature in the space between the graphite plate and the ingot top were also measured. Three ingots A, B and C were cast. Ingot A was solidified in the manner mentioned above. Just before the top of ingot B reached a pasty state, A1-CuA12 eutectic alloy melt held at a temperature of 10 K below the liquidus tem-
Macrosegregation and the a m o u n t of porosity were numerically calculated for the ingots obtained in the present work, which consisted of columnar dendrites growing upwards from the bottom. The method of calculation was the same as that described by Kubota et al. [ 3 ] except that the formation of porosity was considered in the present work. For the purpose, an additional equation to describe the formation of porosity was used, which held in the pasty zone:
V~AFv = -- aV~ A fL -- ~, fLVSj At
(1)
J
The flow rate U in the pasty zone can therefore be obtained from the following equation.
Z
i PLfLUSj
AtS1 ~
Ps(h + ~m) CLi(ko--1) (Phi-- Ps)Pq~
-cLi)
(h + pu~
CpL(P~ -- ps)V~(T-- T~)
(Pu-- Ps)(PsfslCps + P~fuCpL) Vim( CL--Cu) ~,
PufLi _ (p=psV a f v
8j(Ti -- T~) At Z J l/a,~- Axdk~- Axdkj (2)
261
PsV~{ =
~'=1+
-
(Psf~Cps+~pLifuCpL)mc (ko-1)j
(h + ~m)psC~(ko-
1)
(4)
g N = 8 . 8 X 10-19dll'3d22"4fL3"2
4. RESULTS
PL~(PI~ - - P s )
Calculation was carried o u t on the basis of the following two models. In the first model (model I) which corresponded to ingots B and C, the capillary pressure was assumed to be negligible. According to this model, the interdendritic liquid at the t o p of the original b o d y of the ingot flows down until solidification of the region is completed. In the second model {model II) which corresponded to ingot A, it was assumed that the interdendritic liquid stopped flowing because of the capillary pressure after the fraction solid at the ingot t o p exceeded 0.2. This assumption resulted from the observation by O k a m o t o [4] that secondary arms growing from neighbouring columnar dendrites began to t o u c h each other at a solid fraction between 0.16 and 0.22. The following initial and b o u n d a r y conditions were adopted. At time zero in the bulk liquid, the temperature of the melt is uniform and equal to the pouring temperature and the melt is quiescent. The initial temperature of the mould is 353 K and the ambient temperature of the mould is 298 K. The flow rate is zero on the mould wall. The temperature in the space b e t w e e n the graphite plate and the ingot top is 1123 K. The physical and thermal properties of A14.5wt.%Cu alloy, mould and aluminium chill are given in Table 1 and in Appendix A. The permeabilities Kp and K s for liquid flow perpendicular to the columnar dendrites and normal to the columnar dendrites respectively are as follows [ 5 ] : K p : 6.2 X lO-3d12"2d2-1"7fL3"2
(3)
The measured cooling curves in ingot A at three locations 5, 55 and 105 mm from the chill face are shown by the full lines in Fig. 3. The d o t t e d lines in the figure are the cooling curves calculated on the basis of model II and are in good agreement with the experimental lines. Figure 4 shows the measured cooling curves and the calculated curves for ingot B. It is apparent from Fig. 3 or Fig. 4 that there is a positive temperature gradient from the b o t t o m to the t o p of the ingots during the whole period of solidification. The arrow in Fig. 4 shows the time at which A1-CuA12 eutectic melt was added onto the ingot top. Only for a short period after the addition did the temperature slightly decrease at the location of thermocouple 3 and in the upper portions. The macrostructures in the longitu-
1073 /
i
973
,,=,
873
~
........
•
5 .....
........ 1
773
673
, 300
i 600
, 900
, 1200
, 1500
1800
TIME ( s )
Fig. 3. Measured ( ) and calculated (..... ) cooling curves in ingot A. The numerals 1, 3 and 5 represent the thermocouples in Fig. 2. T e is the eutectic temperature of the AI-4.5wt.%Cu alloy.
TABLE 1 Values of the density, thermal conductivity and specific heat used in the calculation
AI-4.5wt.%Cu, liquid AI-4.5wt.%Cu, solid AI-4.5wt.%Cu, eutectic solid Brick Aluminium chill
p
~
ce
(kg m -3)
(W m -1 K -1)
(kJ kg -1 K -1)
PL = PO + h(Cb -- Co) + ~(T-- TO)
138.1 138.1 138.1 138.1 138.1
1.088 0.896 0.896 0.837 1.075
2620 3400 670 2700
262
dinal sections of ingots A, B and C are shown in Fig. 5. The densities of the porosity free hot isostatically pressed specimens are plotted against the copper content in Fig. 6. This relation is formulated as follows: p = 2700 + 22 × wt.% Cu
(5)
where p is the density. The amount of porosity in the ingots was obtained from the density measurement, as is shown in Fig. 7. The data for ingots B and C are neglected near the ingot tops because
the added low melting point liquid metal penetrated into the ingot over some distance from the tops. It can be seen that the addition of the low melting point liquid metal onto the ingot top resulted in a significant decrement in the amount of porosity compared with the case of no addition. The amount of porosity calculated on the basis of model II, which corresponds to ingot A, becomes larger with increasing distance from the chill surface as shown by the broken line in Fig. 7. The amount of porosity calculated from model I corresponding to ingots B and C is essentially zero throughout the ingots. The calculated
1073 3.0 xl03
A
973 2.9
I 873
/I
28
,,=,
Te = 821 K
t-
~
o ~~
I +o ~ + + ~ +
,
o++
++
773 2.7, 673 0
i 300
tPl, +
900
600 TIME
L
i
1200
1500
1800
( s )
Fig. 4. Measured ( ) and calculated ( . . . . . ) cooling curves in ingot B. tp is the time at which AI-Cu eutectic ahoy liquid was added onto the ingot top.
2,6 0 0
' 1.0
' 2.0
' 3.0 Cu
' 4.0
' 5D
CONTENT (
wt
' 6.0
7.L0
' &O
' 9.0
150
pct)
Fig. 6. Densities of AI-Cu alloys free from porosity: o, measured; e, ref. 6.
mm
100
50
Fig. 5. Macrostructures on the longitudinal sections of unidirectionally solidified AI-4.5wt.%Cu alloy ingots: (a) ingot A; (b) ingot B; (c) ingot C.
263
4.0
3.0
e°
3 F--
o no Q.
2.0
::::t 14.0[-
I---
zt ~
6.0
b--
z
ot 9
1.0
5.0 ~-%~- ~ ~,_ _---_--__ "__ _ ~ * _ ~ _
...........
3
c)
0.0
O
I
I
I
I
I
20
40
60
80
100
120
4.0
3.0
DISTANCE FROM CHILL ( m m )
0
2'0
40
I 60
, 100
J 80
120
DISTANCE FROM C H I L L ( m m )
Fig. 7. Amount of porosity in ingots A (o), B (e) and C (c~): - - -, value calculated from model II.
Fig. 9. Macrosegregation in ingot B: o, experimental results;--, results calculated f r o m m o d e l I; - - -, results calculated f r o m m o d e l II.
5.5 o
~" ,...;
18.0 /
5.0
17.0 II-Z
,,,
/-,.5
I,Z
o (.9
v FZ
4.0 3.5
LU
13.OJ~
ou
T 60 r
)-Z
0
210
41O
, 60
I 80
= 100
120
DISTANCE FROM CHILL ( m m )
Fig. 8. Macrosegregation in ingot A: e, experimental r e s u l t s ; - - , results calculated from model I; - - -, results calculated from model II.
a m o u n t o f p o r o s i t y in t h e ingots is smaller t h a n t h e e x p e r i m e n t a l value. T h e d i f f e r e n c e s between the calculated and the measured values f o r t h e t h r e e ingots agree w i t h e a c h other. T h e m e a s u r e d m a c r o s e g r e g a t i o n in i n g o t A is s h o w n b y closed circles in Fig. 8. T h e c o p p e r c o n c e n t r a t i o n d e c r e a s e s w i t h increasing dist a n c e f r o m t h e chill s u r f a c e , i.e. inverse segreg a t i o n occurs. Also in t h e figure are s h o w n t h e c a l c u l a t e d results f r o m m o d e l s I a n d II. T h e m a c r o s e g r e g a t i o n is larger in m o d e l I t h a n in m o d e l II. T h e d e g r e e o f m a c r o s e g r e g a t i o n c a l c u l a t e d f r o m m o d e l I I agrees q u i t e well w i t h t h e e x p e r i m e n t a l values. T h e m a c r o s e g r e g a t i o n in i n g o t B is s h o w n in Fig. 9. I n this case, g o o d a g r e e m e n t is o b t a i n e d b e t w e e n e x p e r i m e n t a n d c a l c u l a t i o n o n t h e basis o f m o d e l I. I n Fig. 10 t h e c o p p e r a n d lead c o n -
14.0~ °
•
•
°
s.o F ~ - ~ . . _ . , _ ........ - - ~ _ : : : : : : U
..........
40k
:o/0
,
,
,
20
40
60
80
100
120
DISTANCE FROM CHILL ( r a m )
Fig. 10. Distribution of copper and lead in ingot C: o, experimental results for copper; o, experimental results for lead; - - . , results calculated from model I; - - -, results calculated from model II.
t e n t s in i n g o t C are p l o t t e d . I t can be seen t h a t lead p e n e t r a t e d i n t o t h e ingot. 5. DISCUSSION T h e solidification o f all t h e ingots in t h e present work proceeded upwards from the chill f a c e . T h e liquid f l o w t o c o m p e n s a t e f o r solidification s h r i n k a g e in t h e ingot, t h e r e f o r e , s h o u l d o c c u r vertically d o w n w a r d s f r o m t h e i n g o t t o p . I n i n g o t A, t o w h i c h n o l o w m e l t i n g p o i n t liquid m e t a l w a s a d d e d , t h e f e e d i n g process c a n be divided i n t o t w o stages. In t h e first stage, t h e r e is a b u l k liquid r e g i o n a b o v e t h e
264
pasty zone; the bulk liquid is supplied into the pasty region to feed solidification shrinkage in this region. In this case the liquid flows easily against liquid flow resistance coming only from viscous friction in the pasty zone. In the second stage where the ingot top becomes in a pasty state, the capillary pressure appears because of the surface tension of the interdendritic liquid and acts as an additional resistance to the interdendritic liquid flow, accelerating the formation of shrinkage porosity and blow holes. This situation is also explained as follows. The pressure of the interdendritic liquid just beneath the free surface is smaller (by the capillary pressure) than the atmospheric pressure. The pressure of the liquid in the ingot is therefore much smaller than that in the case where there is no capillary pressure. As a result, pores are easily formed in the ingot. The t o p surface of ingot A was rather smooth after solidification had been completed, and there was no dendrite arm protruding from the top surface to the air. This fact indicates that the capillary pressure is very large and that no interdendritic liquid could flow down in the second stage in ingot A. The results in Fig. 7 show that the a m o u n t of porosity decreases largely by the addition of low melting point liquid alloy onto the ingot top. The liquidus temperature of the top layer of ingot B was lower than that of A1-4.5wt.%Cu alloy because of the addition of eutectic alloy melt. The capillary pressure, therefore, began to arise in a much later stage of the solidification than in the case of ingot A. At that time, the fraction of liquid in the original part of the ingot was small and the amount of porosity should be small. The lead layer at the top of ingot C, shown in Fig. 5(c), was liquid even when the temperature at the ingot top fell below the eutectic temperature of the A1-CuA12 eutectic alloy system. In this case, the capillary pressure arises because of the interfacial tension b e t w e e n the interdendritic alloy liquid and liquid lead, which is much smaller than that from the surface tension. Feeding liquid from the ingot top is therefore much easier in ingots B and C than in ingot A. This suggests that a large portion of the porosity in ingot A was generated after the t o p of the ingots solidified to become a pasty zone. The fact that the amount of porosity calculated as the basis of model I
or II for each ingot is smaller than the experimental value may be ascribed to the evolution of hydrogen gas during solidification of the ingots. Inverse segregation arises because the soluterich interdendritic liquid flows towards the chill surface to compensate for the solidification shrinkage [7, 8]. The capillary pressure appearing at the t o p of ingot A makes the flow of solute~enriched interdendritic liquid more difficult and decreases the macrosegregation. The resultant segregation in ingot A agrees quite well with the calculated value in which the liquid flow rate is assumed to be zero after the ingot t o p had become in a pasty state. In contrast, when low melting point liquid alloy was added onto the t o p of ingot B, feeding to the ingot was made relatively easy. The segregation can be successfully predicted on the assumption of complete flow-back of soluteenriched interdendritic liquid as shown in Figs. 9 and 10. It can be said from the results in Figs. 8 - 1 0 that, the larger the amount of porosity, the smaller the degree of segregation is.
An oxide film is easily formed on the surface of liquid aluminium in air. In fact, the top surface of ingot A was covered with an oxide film. Kahl and F r o m m [9] showed that the strength of the oxide film on the liquid surface of pure aluminium and aluminium alloys is much larger than the surface tension of the corresponding liquid. In such a case, after the ingot top becomes in a pasty state, the interdendritic liquid flow would be much more difficult than in the case where the liquid-free surface is free from an oxide film. It is concluded that, when the free surface of a riser or risers becomes in a pasty state, since the capillary pressure or the surface oxide film makes the feeding of liquid to ingots and castings difficult, it accelerates the formation of porosity and also reduces the macrosegregation, in the case of alloys with a wide freezing range.
6. CONCLUSIONS
A1-4.5wt.%Cu alloy ingots were solidified unidirectionally upwards and, just before the ingot t o p became in a pasty state, A1-CuA12 eutectic alloy melt or lead melt was added onto the t o p to keep the t o p region in a bulk
265 liquid state during solidification of the original part of the ingot. When the ingot top becomes in a pasty state, the free surface of the liquid descends into the interdendritic space and causes a capillary pressure to develop which increases the resistance to interdendritic liquid flow. The addition of the low melting point liquid keeps the resistance to interdendritic liquid flow at the same level as that from only viscous friction, because it prevents the occurrence of capillary pressure in the liquid. The interdendritic liquid can therefore easily move downwards into the pasty region to compensate for the solidification shrinkage. As a result, the degree of macrosegregation becomes larger and the a m o u n t of porosity becomes smaller in these ingots than in an ingot to which no liquid metal was added. The macrosegregation in a melt~added ingot can be successfully predicted by computer simulation on the assumption that there is no resistance to the interdendritic liquid flow at the ingot top, whereas t h a t in a no-melt-added ingot can be predicted by assuming t h a t the interdendritic liquid stops flowing after the fraction solid at the ingot top exceeds 0.2. The difference in the a m o u n t of porosity between the ingots with and w i t h o u t the melt addition is qualitatively predicted.
4 T. Okamoto, Tetsu To Hagan$, 58 (1982) 1302. 5 K. Murakami, S. Shiraishi and T. Okamoto, Acta Metall., 32 (1984) 1423. 6 E. A. Brandes, Smithell's Metals Reference Book, 6th edn., Butterworths, London, 1983, p. 14-14. 7 J. S. Kirkaldy and W. V. Youdelis, Trans. MetaU. Soc. AIME, 212 (1958) 833. 8 M. C. Flemings and G. E. Nereo, Trans. Metall. Soc. AIME, 239 (1967) 1448. 9 W, Kahl and E. Fromm, Z. Metallkd., 75 (1984) 957.
APPENDIX A: NOMENCLATURE CT.
solute concentration of liquid
Cp~. specific heat of liquid
This work was supported by the research project of Institute of Scientific and Industrial Research, Osaka, on the development of new materials for energy. The authors wish to t h a n k Professor M. Koizumi, Professor Y. Miyamoto and Dr. K. Suganuma for their help in hot isostatic pressing, and Mr. M. Ona for performing the experiments.
Cps specific heat of solid Co = 4.5 wt.%, solute concentration of liquid at the liquidus temperature dl primary arm spacing d2 -- 50 m, secondary arm spacing fL volume fraction of liquid fs volume fraction of solid Fv a m o u n t of porosity h = 26.3 m -s (wt.%) -1, constant in Table 1 h a heat transfer coefficient H -- 397.5 kJ kg-1, heat of fusion k 0 = 0.172, equilibrium distribution coefficient K N permeability normal to columnar dendrites Kp permeability parallel to columnar dendrites m = -- 3.39 K (wt.%y 1, liquidus slope S area of the side of an element t time T temperature To = 918 K, liquidus temperature Ax distance from a nodal point to the grid line U flow rate V volume of an element
REFERENCES
a
1 C. Y. Liu, K. Murakami and T. Okamoto, Acta Metall., 34 (1986) 159. 2 C. Y. Liu, K. Murakami and T. Okamoto, Acta Metall., 34 (1986) 1173. 3 K. Kubota, K. Murakami and T. Okamoto, Mater. Sci. Eng., 76 (1986) 67.
PL Ps P0
ACKNOWLEDGMENTS
Greek symbols
change in volume on solidification = -- 0.3 kg m -3 K -1, constant in Table 1 thermal conductivity density of liquid density of solid = 2440 kg m -a, density of liquid at the liquidus temperature