Favorable Environment for a Nondendritic Morphology in Controlled Diffusion Solidification
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- Khalaf, A.A. & Shankar, S. Metall and Mat Trans A (2011) 42: 2456. doi:10.1007/s11661-011-0641-z
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The novelty of the controlled diffusion solidification (CDS) process is the mixing of two precursor alloys with different thermal masses to obtain the resultant desired alloy, which is subsequently cast into a near-net-shaped product. The critical event in the CDS process is the ability to generate a favorable environment during the mixing of the two precursor alloys to enable a well-distributed and copious nucleation event of the primary Al phase leading to a nondendritic morphology in the cast part. The turbulence dissipation energy coupled with the undercooling of the precursor alloy with the higher temperature enables the copious nucleation events, which are well distributed in the resultant mixture.
- Alloy 1
precursor alloy with higher thermal mass (higher temperature and higher mass)
- Alloy 2
precursor alloy with lower thermal mass
- Alloy 3
resultant desired alloy
- C1 and C2
average solute concentration of Alloys 1 and 2, respectively
transient solute concentration (Cu)
- m1, m2, and mT
mass of Alloy 1, Alloy 2, and total mass, respectively
- T1 and T2
temperatures of Alloys 1 and 2, respectively
temperature of point B
- TL1, TL2, TL3
liquidus temperatures of Alloys 1, 2, and 3, respectively
transient instantaneous position
A nondendritic morphology of the primary phase in a solidified binary alloy could be achieved by controlling the solute redistribution and thermal fields in the solidifying phases both individually and together. The first such efforts were in the rapid solidification process wherein a nondendritic (cellular) microstructure could be formed in the solidified part, because the extraordinarily high growth rates experienced by the solid-liquid interface lead to a stable growth of the same. In the latter part of the twentieth century, the thixoforming and rheocasting processes were developed as casting methods wherein a nondendritic morphology of the primary phase could be obtained in the cast part with a binary eutectic alloy along with the benefits in improved mechanical properties and performance. Unfortunately, these processes have proven prohibitive in capital and manufacturing costs.[2–7] Aluminum alloys have been the most significantly researched material to obtain a nondendritic cast microstructure, and near-net-shaped casting of Al wrought alloys along with their superior properties and performance have been a challenge for conventional casting routes due to the main disadvantage of hot tearing or hot cracking during solidification as caused by the presence of a pronounced dendritic network, which renders the cast components ineffective. To circumvent the disadvantages of thixoforming and rheocasting processes and to enable a cost-effective near-net-shape casting of Al wrought alloys, the controlled diffusion solidification (CDS) process was innovated to enable casting aluminum alloys with a nondendritic morphology of the primary Al phase in the resultant cast microstructure by circumventing the problem of hot tearing and obtaining a product with improved mechanical properties.[9,10]
In Figure 1, Alloy 1 would be pure Al, Alloy 2 would be Al-33 wt pct Cu, and Alloy 3 would be the desired Al-4.7 wt pct Cu alloys.
Controlled CDS experiments were carried out in a laboratory setting such that two precursor alloys could be mixed and allowed to solidify in a crucible. Two rates of solidification were carried out: a slower rate wherein the mixed resultant alloy was allowed to solidify in a crucible placed on a table and a higher rate of solidification wherein the mixture of resultant alloy was quenched on a rotating copper wheel in ribbon casting equipment so as to interrupt the solidification process at a predetermined time-step. The purpose of using an extraordinarily high growth rate with the ribbon casting equipment was to effect a nearly complete solute trapping, leading to a diffusionless transformation, thereby arresting the solidification process at the predetermined time-steps to investigate nucleation and growth of the primary Al phase.[13,14]
Experiment Design with Independent Parameters and Constants for the Laboratory CDS Experiments
Rate of Mixing (g/s)
T1 [K (°C)]
TL1 [K (°C)]
T2 [K (°C)]
TL2 [K (°C)]
remelt Alloy 3 to 936.15 K (663 °C) from CDS2 and pour into empty crucible at 828.15 K (555 °C) to solidify
solidify Al-4.7 wt pct Cu alloy from 953.15 K (680 °C) in a crucible
The solidified samples were suitably sectioned and mounted using black phenolic (Black Phenolic, MetLab Corporation, Niagara Falls, Canada) and were ground and polished in automatic polishing machines. Etching was carried out on some samples with freshly prepared Keller’s reagent (1 mL HF, 1.5 mL HCL, 2.5 mL HNO3, and remainder H2O). The microstructure analysis was carried out using the stereo, light optical, and scanning electron microscopes. A stereo microscope type NIKON AZ 100 M3 was used along with image analysis software.
To identify the nucleation events during the mixing process, the resultant mixtures of the CDS2, CDS3, and CONV2 experiments shown in Table I were quenched at about 104 °C/s in ribbon casting equipment. The ribbon casting process equipment shown in Figure 2 was used in this study to quench the CDS mixture at a very high quenching rate so as to study the nucleation of the primary Al phase during various stages of the CDS process.
- (1)The copper wheel shown in Figure 3(b) is used to quench the mixture. The wheel was rotated at 1100 rpm in all the quenching experiments.
A long steel tube (423 mm) of 21.12-mm inner diameter was fixed with one end at about 3 mm above the copper wheel and the other end coming out at the top of the chamber, as shown in Figure 3(c). The end of the tube closest to the Cu wheel was closed with steel, and a 5-mm-diameter orifice was drilled in the center to enable the flow of the mixture onto the wheel. A ceramic funnel with a 7-mm diameter of outlet orifice was fitted on the top end of the steel tube, as shown in Figure 3(c), and the mixture was poured into this funnel. The total time between pouring the mixture into the funnel to the first liquid in contact with the wheel was about 0.3 seconds.
An induction copper coil shown in Figure 3(b) was used to heat the bottom of the steel tube to 973.15 K (700 °C) to prevent any appreciable heat loss from the melt mixture.
Experimental Design with Designation, Independent Parameters, and Constants for the Interrupted Quenching Experiments with the Ribbon Casting Process
Time After Mixing (s)
T1 [K ( °C)]
TL1 [K ( °C)]
T2 [K ( °C)]
TL2 [K ( °C)]
Cu (wt pct)
3 Results and Discussion
The value of NWe was evaluated as 30, which is far greater than unity, suggesting that the inertial forces during mixing are far greater than the surface tension forces, which would result in a significant breakdown of the stream of Alloy 1 entering the mixture and distributing the Alloy 1 as small pockets of liquid. Figure 5(a) shows that at any time during the mixing process, the temperature of the resultant mixture would be significantly lower than the liquidus temperature of Alloy 1; further, since the time of mixing is merely a couple of seconds long, the inter-diffusion of the species between Alloys 1 and 2 would be insignificant. Hence, the small pockets of Alloy 1 would experience a large undercooled environment leading to a nucleation event of the Al phase from the various pockets of liquid Alloy 1. Since these pockets of liquid are numerous and well distributed in the resultant mixture, there would be copious nucleation events in the resultant mixture during mixing and this would lead to a favorable environment for the stable growth of the nuclei leading to a nondendritic morphology of the primary Al phase in the resultant casting.
Figure 5(b) shows the temperature of point B, TB, as a function of the T1 wherein with increasing values of T1, TB initially decreases and subsequently shows a gradual increase. The value of TB depends on four parameters: rate of mixing, heat added into the resultant mixture from the enthalpy of Alloy 1, heat added to the mixture by the heat of fusion resulting from the nucleation events of primary Al during mixing, and heat loss to the environment from the resultant mixture. The rate of mixing and heat loss to the environment would be nearly constant for a specific experimental design. Hence, the variation in the value of TB would primarily be due to the variation in the heat added to the mixture during mixing from the enthalpy of Alloy 1 and the nucleation reaction, alike. If there were no nucleation events of the primary Al phase from Alloy 1 in the resultant mixture during the segment AB, TB would continuously increase with increasing values of T1. If there are nucleation events in the resultant mixture during mixing, then enthalpy of fusion would be released into the mixture from these events and, thereby, increase the heat content of the mixture, leading to an increase in the value of TB. The higher the number of nucleation events, the higher would be TB; and the amount of nucleation events would increase when the temperature T1 decreases for a constant value of T2, because the lower superheat of Alloy 1 would result in a higher number of nucleation events for the same rate of heat extraction by Alloy 2 during mixing. In Figure 5(b), the value of TB decreases between the superheat temperatures of 275.15 K and 278.15 K (2 °C and 5 °C) because of the copious nucleation events at the lower superheat, adding heat of fusion to the mixture to result in a higher value of TB. After the superheat of 278.15 K (5 °C), the value of TB steadily increases due to the high enthalpy of Alloy 1 with increasing superheat temperatures and nominal nucleation events.
The results shown in Figures 4 through 10 demonstrated the existence of a unique and favorable environment for copious and well-distributed nucleation events in the resultant mixture of the desired alloy composition caused by mixing an alloy with a higher thermal mass into one with a lower thermal mass. The copious nucleation formed during and after the mixing process along with the thermal and solute redistribution would make a favorable environment around the growing nuclei to increase the stability of the solid liquid interface and result in a predominantly nondendritic microstructure, which is favorable to improve the mechanical properties due to the decrease of the hot tearing and porosity defects associated with dendritic morphology in conventiontory experiments. The critical condition for copious nucleation taking place is by setting up the first precursor alloy (Alloy 1) temperature closer to its respective liquidus temperature.
A simple analogous example could be that of making coffee wherein if the higher thermal mass of black coffee is mixed into a lower thermal mass of cream at nearly room temperature, one would observe a nearly uniform mixture of coffee without the aid of stirring, whereas if the colder cream is mixed into the hotter coffee, stirring of the coffee is mandatory to obtain a nearly uniform mixture.
The observations of this study could also be used to develop novel methods of obtaining semisolid slurries in slurry ready semisolid metal casting processes such as the new rheocasting process, continuous rheoconversion process, subliquidus casting process, semisolid rheocasting process, and rheo die-casting process. In all these semisolid casting processes, a large amount of capital and processing costs are required to obtain nearly homogeneous thermal and solute fields in the semisolid slurry mixture prior to the casting process. CDS would enable such a mixture with minimal cost and use the nature of liquid mixing to obtain a favorable semisolid slurry material.
The authors acknowledge the financial support provided by the Natural Science and Engineering Research Council (NSERC) of Canada through their Discovery Grant Program.