The Effect of Copper on Feeding Characteristics in Al-Si Alloys

Abstract

The effects of Cu on feeding and macro-porosity characteristics were investigated in hypo- (A356 and 319) and hypereutectic (391) aluminum-silicon alloys. T-section and Tatur tests showed that the feeding and macro-porosity characteristics were significantly different between the hypo- and hypereutectic alloys. The hole and the pipe in the T-section and the Tatur casting in hypereutectic alloy showed a rough and irregular shape due to the faceted growth of the primary silicon, while the results of the hypoeutectic alloys exhibited a rather smooth surface. However, the addition of Cu did not strongly affect the macro-feeding behavior. It is known that copper segregates and interferes the feeding process in the last stage of solidification, possibly leading to form more amount of micro shrinkage porosity by the addition of Cu. The macro porosity formation mechanism and feeding properties were discussed upon T-section and Tatur tests together with an alloying addition.

1. Introduction

Demands for light weighted aluminum casting alloys have been increased for structural applications in automotive, marine and aerospace industries. The aluminum-silicon alloys are known to be one of the most important casting alloys because of their superior casting characteristics. The performances of the silicon-based aluminum alloys are mainly affected by the size and morphology of the aluminum-silicon eutectic phases; therefore, a considerable amount of research has been carried out. At the same time, since copper improves strength and hardness of aluminum alloys by precipitation hardening in both as-cast and heat-treated condition, copper has also been used as one major alloying element in aluminum alloys [1].

However, it has been documented that copper is usually responsible for a reduction in corrosion resistance and hot tearing resistance [1-4]. Also, casting characteristics of the aluminum alloys are generally biased upon a copper addition, because copper and its compounds nucleate and grow in the last stage of solidification and appear to interfere with feed metal transfer, resulting in the formation of an increased amount of shrinkage porosity and microporosity in the alloys [5,6]. Therefore, aluminum alloys containing silicon and/or copper as both a major and minor element limits a practical application in some conditions where soundness of castings is critical and/or the castings are in service in a severe corrosive environment.

Especially, the feeding and micro-porosity characteristics are of significant for fabrication of sound casting products, while there are few documents concerning the effect of silicon and/or copper on the solidification of aluminum alloys. It has been documented that the solidification/ feeding behavior of aluminum-silicon alloys is affected by composition variation [5,7], while the mechanism by which copper interferes with the feed metal transfer has not been fully understood.

In this paper, a special focus has been given to identify the underlying mechanism of the feeding and macroshrinkage formation characteristics affected by aluminumsilicon composition differences and in the variation of copper content as well. Two series of experiments, Tsection casting and the Tatur test, were conducted to observe the copper-modified three 3XX series aluminum casting alloys including A356.2 (Al-7%Si-0.3%Mg; hypoeutectic alloy), 319 (Al-6%Si-3.5%Cu; hypoeutectic alloy) and 391 (Al-18%Si-0.6%Mg; hypereutectic alloy).

2. Experimental Procedures

Two different kinds of hypoeutectic aluminum-silicon alloys, A356 and 319 alloy and one hypereutectic alloy, 391 alloy, were prepared for the investigation. The compositions of each alloy are shown in Table 1. Melts of about 5 kg were prepared in an induction furnace. The crucible was washed with a pure argon gas and a positive pressure of argon was kept throughout the melting procedure by dripping liquid argon on the melt. After the alloys were completely melted, the melts were held for 15 minutes to improve homogeneity at the temperature of 740ºC for the A356 and 319 alloys and 780ºC for the 391 alloy, about 130ºC above the liquidus temperature of each alloy. The calculated liquidus, solidius temperatures and freezing ranges of the alloys studied were shown in Table 2 by using Pandat software. At those holding temperatures, copper in a form of high purity (99.99%) was added into the melts. Copper additions were made 0% and 0.5% for A356 and 391 alloy, and 0% and 1.0% for 319 alloy, resulting that total amounts of copper were 0.01% and 0.51% for A356, 3.36% and 4.36% for 319 alloy, and 0.02% and 0.52% for 391 alloy. By addition of copper, copper contents of high copper alloys were above of the maximum composition limits of each alloy. Following a holding at 780ºCC for the A356 and 319 alloy, and 800ºC for the 391 alloy for 5 minutes, melts were poured into the T-section and Tatur molds. The composition ranges of each alloy system and compositions for cast alloys are shown in Table 3.

Table 1. Compositions of the aluminum alloys (wt.%).

Table 2. Calculated liquidus, solidius temperatures and freezing ranges of each alloy (ºC).

Table 3. Typical composition range of the commercial alloys and cast alloy compositions (wt.%).

* Maximum unless range is given

T-section casting has been performed to examine the feeding characteristics that can be affected by the differences of the copper amount in the alloys. The neck of the T-section castings, which is designed to promote the shrinkage defects, was examined to observe any differences between the low copper and high copper alloys. The dimension of a graphite T-section mold is shown in Fig. 1. The Tatur test was also conducted to observe the shrinkage characteristics of the aluminum alloys. The dimension of the Tatur steel mold is described in Fig. 2. The temperature of the Tatur mold was controlled to 300ºC prior to pouring.

Fig. 1. Schematic figure of T-section mold and dimension; left, front view; right, cross section (unit: mm).

Fig. 2. Schematic figure of steel Tatur mold and Dimension (unit: mm).

3. Results

Fig. 3 shows typical microstructures of T-section castings of A356, 319 and 391 alloys, and Cu-modified A356, 319 and 391 alloys. The microstructures of both the high and low copper alloys had very little or no influence upon the additional of copper 0.5% for A356 alloy (Fig. 3 (a) and (b)). However, for the 319 alloy with 3.36 %Cu and 4.36 %Cu, there was no significant difference observed except the formation of a larger amount of Al2Cu phase in the high copper alloy (Fig. 3 (c) and (d)). The solidification mode of the hypereutectic 391 alloy is quite different from the hypoeutectic A356 and 319 alloys, because of the nucleation of primary silicon particles before the formation of α-aluminum. However, a clear difference in the microstructure was not observed upon copper addition (Fig. 3(e) and (f)). Development of large and blocky shaped primary silicon was clearly observed, and following that formations of α-aluminum and eutectic solidification were also shown. Note that the eutectic around the primary silicon was developed in a divorced manner and formed a shell of α-aluminum around the primary silicon particles. A number of primary silicon phases show the entrapped, or enveloped, pockets within the crystals. The pockets contained either aluminum or modified-fibrous eutectic.

Fig. 3. Optical Micrographs of; (a) A356 alloy with 0.01%Cu, (b) A356 alloy with 0.51%Cu, (c) 319 alloy with 3.36%Cu, (d) 319 alloy with 4.36%Cu, (e) 391 alloy with 0.024%Cu and (f) 391 alloy with 0.524%Cu.

Fig. 4 shows the cross sections of the T-section castings. The T-section casting was designed to have shrinkage defects at the junction of the T, which is the last solidified region according to Chvorinov’s rule [8]. During solidification, feed metal is supplied from the T-junction to both the cross arms and the supporting arm to compensate for the volumetric shrinkage. Therefore, the formation of shrinkage porosity is usually observed at the junction of the T.

Fig. 4. Cross sections of the T-section castings; (a) A356 alloy with 0.01%Cu, (b) A356 alloy with 0.51%Cu, (c) 319 alloy with 3.36%Cu, (d) 319 alloy with 4.36%Cu, (e) 391 alloy with 0.024%Cu, (f) 391 alloy with 0.524%Cu.

The large macro-shrinkage hole was observed at the junction of the T-section casting in A356, 319, and 391 alloys. Moreover, hot tearing (or hot cracking) was also evident on the surface of the T-section castings in A356 and 319 alloys, while no surface defect was observed on the surface of the castings of 391 alloy. Judging from the size or volume of the holes generated in T-section castings, notable discrepancy was not clearly observed among the different amounts of copper. However, it was evident that the high copper A356 alloy exhibited a larger amount of dispersed shrinkage porosity in the cross section of the Tsection casting as shown in Fig. 4 (b). The shrinkage hole geometry produced in the hypereutectic 391 alloy was quite different from the hypoeutectic A356 and 319 alloys because of the interference of liquid feeding during solidification due to the nucleation of large amount of primary silicon before the formation of α-aluminum and its faceted growth mechanism. While the shape of shrinkage porosity of the A356 and 319 alloys were shown to be rather round and smooth, that of the 391 alloy was irregular and rough.

The cross sections of the Tatur test castings are shown in Fig. 5. The pipe volume represents the amount of shrinkage in the Tatur test. Because of the cone shaped mold, the molten metal is continuously fed from the top to the middle of the casting, which is the last solidified region, and generates the pipe, which is considered as the total amount of macro-shrinkage in the alloy. Again, because of the faceted growth mechanism of hypereutectic alloys, the pipe geometry was shown to have a rougher surface and more irregular shape. Fig. 6 shows the pipe volume and total shrinkage area measured from the Tatur castings. Image analysis was used to measure the total macro-shrinkage area. It is evident that the differences in the amount of copper did not markedly influence the macro-shrinkage formation in each alloy. However, it is manifested that the shape of the pipe of the high copper 391 alloy (0.524%Cu) was curved pipe geometry, while that of the low copper 391 alloy (0.024%Cu) was shown to be rather straight, indicating that the feeding mechanism is affected by the presence of a higher amount of copper in the 391 alloy.

Fig. 5. Cross sections of the Tartar castings; (a) A356 alloy with 0.01%Cu, (b) A356 alloy with 0.51%Cu, (c) 319 alloy with 3.36%Cu, (d) 319 alloy with 4.36%Cu, (e) 391 alloy with 0.024%Cu, (f) 391 alloy with 0.524%Cu.

Fig. 6. Pipe volume measured from the Tatur test castings.

4. Discussion

Five feeding mechanisms during solidification of a casting have been described by Campbell [9]; liquid, mass, interdendritic, burst and solid feeding. Among them, interdendritic feeding has been regarded as the most important feeding mechanism for aluminum casting alloys, because they solidify over a wide freezing range and develops extensive mushy zone. Interdendritic feeding describes the flow of residual liquid through the mushy zone. Therefore, it takes place after the formation of the dendritic network. The presence of eutectic like the aluminum-silicon alloys makes the interdendritic feeding easier because the eutectic solidifies at a specific temperature and, therefore, at a specific planar front.

As shown in Fig. 4, both a slump (shown as a hole) and severe hot tearing were observed in A356 and 319 alloys at the surface of the T-section castings, while there were no surface defects exhibited in 391 alloys because of the higher fluidity and lower volumetric contraction of hypereutectic alloys. Hot tearing or hot cracking generally occurs during solidification of casting and welding because of a loss of strength and ductility, resulting from the thermal stress and the presence of intergranular liquid films [10-15]. Due to the volumetric contraction during the transition from liquid to solid (mainly aluminum which contracts up to 8.14% [5]) and no fillet radius in T-section design, both A356 and 319 alloy experienced severe hot tearing.

Because of the differences in the growth mechanism between hypoeutectic and hypereutectic alloys, A356 and 319 alloys showed a rather smooth surfaced hole and pipe in the T-section casting and the Tatur casting. However, the hole and pipe surface of 391 alloy was observed to be rough and irregularly shaped because of the faceted growth mechanism of the silicon phase. Detailed observations of the size and the geometry of the shrinkage holes generated during solidification in the cross section of the T-section castings lead that the amount of copper in each alloy system had little influence on the formation of the macroshrinkage. Since the solidification characteristics were not significant with an increased amount of copper, it is probable that the effect of copper on the macro-shrinkage formation in the 319 alloy is not significant. However, the reason why there was little influence on the formation of macro-shrinkage with the amounts of copper in the A356 and 391 alloy is probably due to the fact that copper interferes with the feed metal transfer mainly after more than 98 or 99% of solidification has been completed. In other words, the micro-segregation effect of copper becomes predominant during the last stage of interdendritic feeding, when the majority of casting has been already solidified. Therefore, the influence of the amount of copper on the formation of the macro-shrinkage hole for a given alloy was not notable.

The Tatur test showed a similar phenomenon as observed in the T-section castings. The differences in the pipe volume represent the total amount of macro-shrinkage and the results illustrated that the pipe volume was not significantly affected by the differences in copper content for A356, 319 and 391 alloys as shown in Fig. 6. The volume were compared by filling sand into the pipe and each amount of the sand was measured. However, the pipe geometry of the 391 alloy showed an evidence of interference on the feeding characteristics of the alloy, as shown in Fig. 5 (e) and (f). While the feed metal was transported rather uniformly to the last solidifying region to form a shrinkage pipe in the low copper alloy (showing a rather straight pipe geometry), an additional 0.5% content of copper made the pipe narrower and more curved. This is probably due to the formation of copper phases at the eutectic cell boundaries and the consequent inhibition of the feed metal transfer during the forming of the shrinkage pipe.

The results described above concerning the interference in the feeding characteristics resulting from the presence of copper in aluminum casting alloys could be explained with several important mechanisms. Due to the extended solidification temperature range, more than 40ºC for both A356 and 391 alloys, and increased eutectic mushy zone area developed by an increased amount of copper in A356 and 391 alloy, the feed metal channels can be greatly extended, thereby making it more difficult for feed metal to flow through the longer interdendritic channel. In addition, the formation of copper-enriched liquid pools at the last solidifying regions will encounter blockage of feed metal transport, because those high copper liquid pools become enveloped by solid eutectic phases, mainly copper phases, as temperature decreases, which will choke or clog the feed metal channels. This is also accompanied by a decrease in the silicon content of the last liquid to solidify in the high copper alloys. While detailed study for scrutinizing macro and micro porosity formation associated with microsegregation is under way, current observations and experiments clearly identify that the feeding behavior and macro-porosity formation is affected by mainly silicon rather than copper.