2025 Volume 66 Issue 1 Pages 93-98
Aluminum casting alloys are widely used for various machine parts for the purpose of weight reduction and energy saving. Because of the strong demands for high-stiffness Al alloy to improve their operating accuracy in the fields of precision machinery and robots, we have developed Al-12.5∼14%Si-5∼6%Ni-14∼15%Cu-0.5%Mg alloys with a stiffness of 100 GPa. The solidification sequence of this alloy was investigated to clarify the structural constituents that contribute to its higher stiffness. A series of quenched Al-14%Si-5.5%Ni-15%Cu-0.5%Mg samples were analyzed by optical microscope, EPMA and X-ray diffraction. The results revealed that the alloy solidifies in the order of primary Si, primary-like Ni2Al3, Al(α) + Si eutectic, Al(α) + Ni2Al3 eutectic, Al(α) + CuAl2 eutectic and Al(α) + Al4Cu2Mg7Si8 eutectic. Just after the crystallization of primary-like Ni2Al3, Al(α) hollow and α-dendrites developed around primary Si and Ni2Al3, since these primaries hardly nucleated the above eutectics. The thermodynamic calculation software used provides information on the solidification temperatures and volume fraction for primary Si, Ni2Al3, and other eutectic structures. However, it is unable to predict the appearance of Al(α) hollow and α-dendrites because these are non-equilibrium phenomena. In addition, some discrepancies were observed especially in the final stage of solidification. The evaluation performed by the thermodynamic calculation software using the Scheil’s equation provides details on the final solidification reactions similar to experiment findings.
This Paper was Originally Published in Japanese in J. JFS 95 (2023) 419–425. Some figure captions were modified.
Aluminum casting product has a higher specific strength than cast iron, excellent corrosion resistance and high thermal conductivity, and good castability. Therefore, in order to reduce the weight and energy consumption of machines, aluminum product has been widely used in automobiles and various other machine parts as a substitute for cast iron [1]. In recent years, many parts that move at high speed are used in the industrial machinery and robotics fields. Since the performance of the machine is determined by the moving speed and the accuracy at the moving and stopping positions, lightweight and high-stiffness materials are required. For example, the gray cast iron, which has been widely used in the mechanical field, has a stiffness of over 100 GPa. On the contrary, the stiffness of aluminum casting alloys, such as the Al-Si-based JIS AC3A alloy and the Al-Si-Mg-based JIS AC4C alloy are around 72 to 77 GPa, which are considerably lower than that of gray cast iron. Even for an Al-Si-Cu-Mg-Ni alloy, which has been standardized as AC9A with a significantly increased Si content of 22–24 mass% (hereafter abbreviated as %) and a large amount of primary Si, its stiffness is only about 90 GPa [2–5]. However, the specific stiffness of aluminum alloys can have higher potential due to their low density [6, 7], so machine manufacturers have been demanding further improvement [8, 9]. Composite materials, which contain dispersed reinforcing materials such as ceramic fibers and particles, have been developed as one of the solutions to significantly improve the specific stiffness of aluminum alloys. However, the ceramic material itself is expensive, the wettability between ceramics and molten metal is insufficient, and the technology for separating and recycling the reinforcing materials after use has not been fully established, making it difficult to apply them to large products in the industrial machinery and robotics fields [10, 11].
In order to develop an alloy with higher specific stiffness than conventional aluminum casting alloys, we focused on Si, Ni, and Cu as alloying elements and have improved stiffness by dispersing and crystallizing various intermetallic compounds containing these elements as primary crystals and eutectic [12, 13]. We have produced numerous sand and permanent mold casting specimens with systematically varying compositions in the ranges of 6–18%Si, 1–8%Ni, and 8–15%Cu, and measured their stiffness and tensile strength. It was found that an Al-12.5–14%Si-5–6%Ni-14–15%Cu alloy exhibits a stiffness of approximately 100 GPa and tensile strength of 80 MPa by a sand mold and 180 MPa by a permanent mold [14–16]. However, this alloy contains primary and eutectic structures with various different morphologies, and many aspects of the contribution of each phase to specific stiffness and the process by which the solidification structure is formed remain unclear. In order to understand the phases that contribute to improved specific stiffness and control their structures, it is first necessary to clarify the solidification process.
In this study, the high-specific-stiffness Al-Si-Ni-Cu-Mg casting alloy containing many primary and eutectic structures was produced, and the solidification structure was clarified by microstructural observation and X-ray diffraction. In addition, the crystallization process of each phase was analyzed by thermal analysis, and the solidification sequence was discussed.
The experimental specimens used in this experiment were prepared by adjusting the composition of 99.5% silicon, nickel, copper, iron, and manganese to the specified composition using a commercially available AC4CH master alloy as the main raw material. The metal ingots were placed in a graphite crucible with an inner diameter of 139 mm and a height of 249 mm, and melted at 1033 K in a gas furnace. The alloy was then cast into a JIS-specified B-type Y-block sand mold and a permanent mold (JIS H 5202) preheated to 473 K. The chemical compositions of the obtained specimen are shown in Table 1. The solidification microstructure of the specimen cast into a B-type Y-block sand mold, which has a relatively slow cooling rate, was examined using an optical microscope, and the compositions of each phase were analyzed using an EPMA (EPMA8050G, Shimadzu Corporation). In addition, X-ray diffraction (D8 DISCOVER, Bruker) was performed using Cu-Kα radiation at 40 kV and 40 mA to identify the crystallized phases.
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To investigate the solidification sequence of this alloy, thermal analysis was performed using specimen, which was cut from permanent mold and has a fine and homogeneous structure. Approximately 5 g of the specimen was placed in a special refractory crucible (53%SiO2-41%Al2O3) and melted by heating to 1033 K in a circular ceramic electric resistance furnace in an Ar gas atmosphere, and then cooled at 0.17 K/s (10 K/min) to obtain the cooling curve. In addition, specimens were prepared by quenching in water immediately after the crystallization temperature on the cooling curve, and microstructural observation and EPMA analysis were performed to evaluate the type, morphology, and element concentration of each microstructure. Meanwhile, the solidification sequence of this alloy was analyzed using integrated thermodynamic calculation software (Thermo-Calc [17], ver. 2022b) with the TCAL8 database, and compared with the experimental results.
The microstructure of the specimen cast into the B-type Y-block sand mold is shown in Fig. 1. Coarse block-like and rod-like structures presumably formed as primary phases, and several types of microstructures presumably formed as eutectic structures, are observed. To clarify the constituent elements of each microstructure, each structure was analyzed by EPMA. The obtained mapping images are shown in Fig. 2. The relative concentration of each element is shown in color, with blue indicating a lower concentration and red indicating a higher concentration. Al is the reference element, and the microstructure with high concentration shown in orange is considered to be the primary α-dendrite phase. In addition, Al is distributed in various concentrations in other primary coarse phases and eutectic structures. The Si-rich phase is determined to be a Si phase because the compositions of other elements are extremely low. And the coarse structure is determined to be primary Si, and the fine structure is determined to be Si crystallized as a eutectic. The Ni-rich rod-like structure, which grows three-dimensionally into a plate shape, contains a large amount of Al and Cu, and it is considered an Al-Ni-Cu intermetallic compound. From a result of X-ray diffraction of this specimen, peaks for many phases were obtained as shown in Fig. 3. By matching the diffraction peaks of Al, Ni, and Cu, a peak corresponding to Ni2Al3 was observed. Therefore, it was determined that the Al-Ni-Cu structure is Ni2Al3 phase with a solid solution of Cu. In Fig. 2, the Fe- and Mn-rich structure also exhibits a coarse structure, and this was considered to be AlFe2Mn by the XRD results in Fig. 3. Fine microstructures presumed to be eutectic solidification are also observed in the analysis maps of Si, Ni, Cu, and Mg in Fig. 2. The structure containing a large amount of Si was considered to be an Al(α) + Si eutectic structure, and the structure rich in Al, Ni, and Cu was considered to be an Al(α) + Ni2Al3 eutectic structure. Furthermore, an Al-Cu-rich and Mg-rich eutectic structures are observed in the analysis maps of Cu and Mg, respectively in Fig. 2. The XRD analysis in Fig. 3 indicates that these microstructures are Al(α) + CuAl2 and Al(α) + Al4Cu2Mg7Si8 eutectic structures, respectively. Representatives of these microstructures are also shown in Fig. 1. Besides, the area fractions of the Si phase and plate-like Ni2Al3 phase, which were observed as primary coarse structures, were 13.7% and 12.9%, respectively. Furthermore, many eutectic structures showed their characteristic morphology. According to Croker’s classification of eutectic morphology [18], eutectic CuAl2 shows a typical lamellar structure, while eutectic Si and eutectic Al4Cu2Mg7Si8 show irregular flakes, and eutectic Ni2Al3 shows anomalous complex regular structure.
Microstructures of specimens of JIS B-type Y-block (a), Enlarged view of JIS B-type Y-block (b) and JIS mold specimen (C). (online color)
EPMA mapping analysis on solidification structure of high-stiffness Al-Si-Ni-Cu-Mg cast alloy. (online color)
X-ray diffraction pattern of high-stiffness Al-Si-Ni-Cu-Mg cast alloy.
To investigate the solidification sequence of the experimental specimen, a thermal analysis curve was initially obtained at a cooling rate of 10 K/min, and is shown in Fig. 4(a). The time differential curve calculated from Fig. 4(a) is also shown in Fig. 4(b). From these two thermal analysis curves, the solidification of this alloy begins at 933 K and completes at 772 K with seven inflection points between 933 K and 770 K, suggesting that seven microstructures were solidified. Therefore, specimens were prepared by quenching the crucible in water from temperatures A to E immediately lower the seven inflection point, and their microstructures were observed.
Thermal analysis curve (a) and deviation of temperature by time of experimental alloy (b). (Cooling Rate: 10 K/min)
Figure 5 shows the newly formed microstructures at each inflection point. In the specimen quenched at temperature A (913 K), only block-shaped silicon phases such as squares are shown, and the rest was a fine solidified microstructure formed by rapid cooling from the melt. Therefore, it is considered that primary silicon phases began to crystallize at 933 K, even though this primary solidification temperature is about 70 K lower than that of the conventional high-stiffness AC9A alloy. In the specimen quenched at temperature B, rod-shaped Ni2Al3 containing Cu as a solid solution was observed in addition to the Si phase, and it is thought that Ni2Al3 phase crystallized from 898 K. Since primary silicon has almost no solid solution with other elements, Ni, Al, and other elements pile up into the surrounding melt during the growth of the silicon, causing them to concentrate and promote the formation of Ni2Al3. The microscopic observations show the area of Ni2Al3 in contact with the primary Si phase is very small, it is considered that the primary Si phase does not have a nucleation effect on Ni2Al3. The microscopic observation shows the area of Ni2Al3 in contact with the primary Si phase is small, indicating that the primary Si phase has no nucleation effect on Ni2Al3. From the second inflection point (898 K), the primary Si phase continues to grow along with the growth of Ni2Al3, and it seems to be a eutectic reaction with a transformation from L (liquid phase) to Si + Ni2Al3. As shown by the arrow ($ \Rightarrow $) in Fig. 5(B), the primary Si phase and the plate-like Ni2Al3 phase continue to grow while in contact with each other. Both Si and Ni2Al3 grow in a faceted regime, and the facet phases in the eutectic structure grows independently of each other [19], so Ni2Al3 is considered to have grown in a primary crystal state like the Si phase.
Microstructures of the specimens quenched at critical stages of solidification at A to E in Fig. 4.
In the specimen quenched from temperature C, the primary-like AlFe2Mn phase is observed, whereas this phase is not present at temperature B, suggesting that it crystallizes at 837 K between temperatures B and C. In addition, Fig. 6 shows the microstructure at the different area in the specimen quenched from temperature C. The dendrite Al(α) phase, small amounts of Al(α) + Si eutectic structure and Al(α) + Ni2Al3 eutectic structure are observed around the primary Si and the primarily-like Ni2Al3. Here, as shown by the dashed circle in Fig. 6 and the EPMA line analysis, small amounts of Si protruding from the primary Si can be seen. This is thought to be part of the Al(α) + Si eutectic structure that has grown directly from the primary Si as a nucleus. Here, as shown by the dashed circle in Fig. 6 and the EPMA line analysis, a small Si protruding from the primary Si is observed. This is thought to be part of the Al(α) + Si eutectic structure that has grown directly from the primary Si as a nucleus. Similarly, Fig. 7 shows the results of EPMA analysis of part of Al(α) + Ni2Al3 eutectic structure that grew directly from the primary-like Ni2Al3. The Ni concentration decreases in locations farther from the primary-like Ni2Al3, while Cu increase. Thus, primary Si and primary-like Ni2Al3 have almost no contribution to the nucleation of the eutectic structure, and it is thought that the Al(α) phase develops into a halo and dendrite shape due to undercooling. This dendritic Al(α) was also observed in the specimens cast in the B-type Y-block sand mold and the permanent mold.
EPMA line analysis across primary Si of specimen quenched at point C in Fig. 4. (online color)
EPMA line analysis across primary-like and eutectic Ni2Al3 in specimen quenched at point C in Fig. 4.
In the quenched structure from temperature D, cellular α + Si eutectic and Al(α) + Ni2Al3 eutectic structures were formed. The quenched specimen at temperature E newly shows the formation of Al(α) + CuAl2 eutectic and Al(α) + Al4Cu2Mg7Si8 eutectic structures, and these eutectics are considered to have crystallized at 779 K and 772 K just before temperature E. From the above observation, it can be summarized that this alloy solidified in the following order: primary Si at 933 K, primary-like Ni2Al3 at 898 K, Al(α) and Al(α) + Si eutectic at 805 K, Al(α) + Ni2Al3 eutectic at 801 K, Al(α) + CuAl2 eutectic at 779 K, and Al(α) + Al4Cu2Mg7Si8 eutectic at 772 K.
3.3 Estimation of solidification sequence using calculated phase diagramTo better understand the solidification process of this alloy, we calculated the relationship between temperature and the volume fraction of each equilibrium phase. The results are shown in Fig. 8. In this alloy, primary Si crystallizes at 933 K, and while primary Ni2Al3 forms at 898 K, the volume fraction of Si also increases. A small amount of Al15Si2(Fe,Mn)4 crystallizes from 875 K. The volumes of the Al(α) phase, Si phase, and Ni2Al3 phase increase rapidly from 819 K, suggesting that the Al(α) + Si eutectic and Al(α) + Ni2Al3 eutectic structures crystallize simultaneously. Furthermore, since the increase in the Al(α) phase from 819 K is significantly greater than the increases in Si and Ni2Al3, it is thought that the volume fraction of the Al(α) phase in the eutectic structure is large.
Change in volume fraction of each solidification phase evaluated at equilibrium condition by Thermo-Calc.
Comparing the experimental solidification sequence shown in Fig. 4 and Fig. 5, the crystallization temperatures of primary Si and primary Ni2Al3 are roughly the same as in the calculated phase diagram. Furthermore, the calculated volume fractions of the primary Si phase and primary Ni2Al3 phase are 8% and 10%, respectively, which are close to the area fractions of 9% and 13% of the B-type Y-block sand cast specimen shown in Fig. 2. However, the composition and crystallization temperature of the primary Al-Fe-Mn compound phase are slightly different. Regarding the eutectic reaction, it is consistent that the Al(α) + Si eutectic and Al(α) + Ni2Al3 eutectic crystallize almost simultaneously, but the eutectic start temperature predicted by the calculated phase diagram is about 15 K higher than the experimental results. The dendrite α phase, which solidifies with eutectic microstructure, cannot be predicted by the calculated phase diagram due to a non-equilibrium solidification phenomenon. Here, the calculation results using the database of all phases are also shown in Fig. 8 with a dashed line. The peritectic reaction of L + Ni2Al3 → Al(α) + Al7Cu4Ni at 816 K is suggested, but since the crystallization of Al7Cu4Ni is not observed in the experimental results, it is determined that this peritectic reaction does not occur. In addition, in the later stage of the eutectic reaction, the equilibrium calculation shows the crystallization of Al(α) + Al5Cu2Mg8Si6, but in the actual solidification, the Al(α) + Al4Cu2Mg7Si8 eutectic structure crystallizes. Although Al5Cu2Mg8Si6 and Al4Cu2Mg7Si8 have slightly different structural formulas, they can be considered to be almost the same phase.
Thermo-Calc has a built-in Scheil’s equation that assumes no diffusion in the solid phase and complete mixing in the liquid phase, so it can calculate the crystallized phase according to the change in concentration of the melt. The relationship between temperature and volume fraction of phases calculated using the Scheil’s equation is shown in Fig. 9. The primary crystal phases are the same as those in the equilibrium calculation, but the eutectic structure formation differs from the equilibrium condition, with Al(α) + CuAl2 and Al(α) + Al5Cu2Mg8Si6 eutectic structures crystallizing at 795 K and 786 K, respectively, these are consistent with the experimental results. The formation of these eutectic structures could be estimated to be caused by changes in solutes in the liquid phase. Therefore, the change in concentration of each element in the residual liquid was also calculated based on Scheil’s equation, and the results are shown in Fig. 10. In the primary crystallization, Si crystallizes at 933 K, decreasing the Si concentration in the liquid phase, and the Ni2Al3 phase crystallizes at 898 K, decreasing Ni in the liquid phase slightly. Regarding the formation of the eutectic structure, the Ni concentration rapidly decreases and the Cu concentration significantly increases with the eutectic solidification of Al(α) + Ni2Al3 at 819 K, that is to say, growth of eutectic Ni2Al3 phase. This corresponds well to the fact that the Ni concentration in the eutectic Ni2Al3 in Fig. 7 decreases the further away from the Ni2Al3 phase the phase is, while the Cu concentration increases. Furthermore, it is thought that the increase in the Cu concentration in this residual liquid promoted the formation of an Al(α) + CuAl2 eutectic structure after Al(α) + Ni2Al3 eutectic solidification.
Change in volume fraction of each solidification phase evaluated using Scheil’s equation by Thermo-Calc.
Evaluation of change in residual liquid composition during solidification by using Scheil’s equation of Thermo-Calc. (online color)
Finally, to clarify the formation process of the solidification structure, a specimen with only the eutectic structure was prepared. From the results of calculations using Thermo-Calc, the composition was determined to be Al-8.8%Si-2.3%Ni-13.0%Cu, when only Al + Si eutectic and Al + Ni2Al3 eutectic structure began to crystallize. Besides, elements such as Mg, Fe, and Mn were excluded. An alloy of this composition was melted and cast into a B-type Y-block sand mold. The microstructure of the specimen is shown in Fig. 11. No coarse primary Si phase or primary Ni2Al3 phase was observed, and although a small amount of primary Al(α) phase was observed, the solidification structure was composed almost entirely of the eutectic structure, suggesting that the formation of the dendritic Al(α) structure seen at 805 K in thermal analysis experiments was due to the crystallization of primary Si and primary Ni2Al3.
Microstructure of Al-8.8%Si-2.3%Ni-13.0% Cu eutectic alloy cast in JIS B-type Y-block.
To obtain fundamental characteristics for the microstructural formation of the high-stiffness aluminum casting alloy with a Young’s modulus of 100 GPa, the solidification process of an Al-Si-Ni-Cu-Mg alloy was analyzed and the following conclusions were obtained.
