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TiAlSi intermetallic formation and its impact on the casting processing in Al–Si alloys
Journal of Materials Processing Technology 210 (2010) 1780–1786
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Journal of Materials Processing Technology journal homepage: www.elsevier.com/locate/jmatprotec
TiAlSi intermetallic formation and its impact on the casting processing in Al–Si alloys X.-G. Chen a,∗ , M. Fortier b a b
University of Québec at Chicoutimi, Chicoutimi, QC, Canada G7H 2B1 Arvida Research and Development Centre, Rio Tinto Alcan Inc., Jonquière, QC, Canada G7S 4K8
a r t i c l e
i n f o
Article history: Received 29 April 2010 Received in revised form 10 June 2010 Accepted 11 June 2010
Keywords: TiAlSi intermetallics Al–Si alloys Microstructure Grain refinement Feeding blockage Wormhole defect
a b s t r a c t The effect of titanium additions on the formation of primary TiAlSi intermetallics and resulted cast microstructure in Al–Si alloys has been studied. The precipitation temperature, growth kinetic and morphology of TiAlSi intermetallics as well as the solubility limit of Ti in Al–Si alloys were experimentally evaluated. The impact of primary TiAlSi intermetallic particles on the casting process related problems was investigated on the large scale of industrial production. The influence of Ti in the occurrence of feeding blockage and wormhole defect during ingot casting was described. The perspectives of Ti addition in Al–Si cast alloys were discussed. It is proposed that the Ti level in the most Al–Si cast alloys should be controlled within 0.10 wt.%. © 2010 Elsevier B.V. All rights reserved.
1. Introduction
aluminum, 665 ◦ C (Mondolfo, 1976):
Aluminum–silicon cast alloys are widely used for the shape casting because of their good castability and mechanical properties. It is a common practice in many aluminum foundries to add titanium in Al–Si cast alloys as a grain refining agent. Although this practice has been employed for many years, Ti addition and grain refining are still a matter of concern to both ingot suppliers and shape casting foundries because of the lack of a definitive knowledge of what is a necessary amount of Ti level to produce adequate grain refining and metallurgical properties. The Ti addition and grain refining practice vary considerable from foundry to foundry, possible ranging from no Ti addition at all to a systematic Ti addition in all cast batches to maintain a uniform in-house operation procedure.
L + TiAl3 ⇔ ˛-Al
1.1. Binary Al–Ti system Before considering the role of Ti in Al–Si alloys, it is useful to first review the binary Al–Ti system. On the aluminum-rich part, Al–Ti alloy presents a peritectic reaction at a composition of approximately 1.2% Ti1 and a temperature above the melting point of pure
∗ Corresponding author at: Department of Applied Sciences, University of Québec at Chicoutimi, Chicoutimi, QC, Canada G7H 2B1. Tel.: +1 418 545 5011x2603; fax: +1 418 545 5012. E-mail address: xgrant [email protected] (X.-G. Chen). 1 Alloy and phase compositions are in wt.% unless specified otherwise. 0924-0136/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2010.06.009
This reaction has a profound influence on nucleation and growth of ␣-aluminum crystals from the liquid stage. TiAl3 intermetallics have an existence range of 36.5–37.5% Ti and a 3370 kg/m3 density (Mondolfo, 1976). The solubility limit of titanium in liquid aluminum is situated between 0.12% and 0.15% Ti (Mondolfo, 1976; McCartney, 1989). Titanium is usually added in the aluminum melt in the form of Al–(5–10%)Ti master alloys. The grain refining is produced by release of numerous TiAl3 intermetallic particles from the master alloys. It is known that these TiAl3 intermetallic particles can act as nucleants for ␣-aluminum crystals during solidification (Arnberg et al., 1982). The formation of TiAl3 intermetallics were studied in the binary Al–Ti system and it was found that TiAl3 intermetallics could have three different morphologies (flakes, petals and blocks) depending on the solidification conditions and the temperature history of the alloy (Arnberg et al., 1982; Guzowski et al., 1987).
1.2. Ternary Al–Si–Ti system At the aluminum-rich corner of ternary Al–Si–Ti system, three possible types of titanium aluminides could be present (Perrot, 1990).
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1. TiAl3 . Up to 15 at.% Al can be replaced by Si in TiAl3 lattice structures, resulting in various chemical compositions and a range of lattice parameters (Mondolfo, 1976). It can be commonly written as Ti(AlSi)3 . 2. 1 : commonly written as Ti7 Al5 Si12 . This phase is stable below 900 ◦ C and can contain 17 up to 42%Si. 3. 2 : commonly written as Ti(AlSi)2 . This phase has higher amount of silicon, which ranges from 38% to 46%.
Table 1 Chemical composition of the investigated alloys (wt.%).
In the most useable Ti ranges of 0–0.3%, all three ternary phases are very close to each other (Perrot, 1990). It can be expected that two or three intermetallic phases can co-exist with aluminum in Al–Si alloys. Because Al and Si can replace each other, it results in a large variety of composition combination and a wide range of lattice parameters of TiAlSi phases. Considering that very little information exists in the literature and equilibriums are hardly achieved in the real casting condition, these titanium aluminides in Al–Si alloys are generally referred to here as TiAlSi intermetallic phases. Habibi et al. (2003) and Fortier and Chen (2004) have recently reported that the morphology of TiAlSi intermetallics in Al–Si foundry alloys is similar to TiAl3 in binary alloy, i.e. flakes and blocks. However, the favourite conditions under which these morphologies would form are still unknown. When Ti concentration added in the melt is less than the solubility limit, TiAlSi intermetallic particles dissolve with time and the grain refining fades away. To avoid this effect, a large amount of Ti (exceeding the solubility) was added for longer lasting refining effect. It is believed that this traditional practice built the basis of Aluminum Association’s chemical composition limits in Al–Si cast alloys many years ago (Sigworth et al., 2007). Although a minimum Ti concentration is not specified, the maximum Ti levels allowed in most Al–Si alloys are very high, mostly in the range of 0.2–0.25%. A number of studies have shown that Ti alone is a relatively weak grain refiner in aluminum alloys in despite of maintaining high Ti level in the melt (Arnberg et al., 1982; Guzowski et al., 1987). The addition of boron and titanium in the form of Al–Ti–B master alloys has been found to be much more effective in grain refinement than titanium alone, since boron in the form of TiB2 enhances significantly the grain refining action. Different Al–Ti–B master alloys with various Ti/B ratios, such as Al5Ti1B and Al3Ti1B, have been developed to improve significantly the grain refinement efficiency in Al–Si cast alloys (Spittle, 2006; Sigworth et al., 2007). An excess of Ti addition can cause a number of problems during liquid metal processing and defects in the final cast products. Sigworth et al. (2007) reported that when a high Ti melt is held in the furnace, TiAlSi crystals precipitated and formed a layer of “sludge” at the bottom of the furnace. Studies of Sokolowski et al. (2000) and Chen and Fortier (2004) revealed that a high level of Ti in the melt created massive TiAlSi crystal precipitation which caused a heavy sedimentation in liquid metal processing lines and a presence of coarse intermetallic crystals in the cast structure. These crystals proved to be detrimental to the foundry processes and the metallurgical properties of the castings. Finally, aluminum foundries obtain their base Al–Si alloys from different sources. The use of pre-inoculated (with Ti) cast ingots is still one important source. In addition, a significant proportion of foundry alloys are in-house recycled materials or mixed with secondary alloys, which usually contained a certain amount of residual Ti. The addition of an excessive Ti combined with non-controlled residual Ti can result in the precipitation of primary TiAlSi intermetallic particles. Therefore, it is of great interest to establish the optimum titanium levels for Al–Si alloys, which can achieve the most desirable metallurgical properties. The aim of the present study is to examine the impact of Ti and its TiAlSi intermetallics on the casting processing in Al–Si cast alloys.
A batch of 5–20 kg of commercial purity aluminum was melted in an electric resistance furnace and alloyed to the desired compositions at 750–770 ◦ C. Three Al–Si cast alloys were studied at different titanium levels ranging from 0.10% to 0.26% and their chemical compositions are given in Table 1. Titanium additions in the melt were achieved using a binary Al–6%Ti master alloy. The evaluation of the formation temperature of titanium aluminides was done using the melt quenching and the LiMCA (Liquid Metal Cleanliness Analyser) detection. In the melt quenching tests, a certain amount of the molten metal (∼100 g) was poured into a small crucible and placed into a furnace at 800 ◦ C. The furnace temperature was gradually lowered to reach the desired temperature. Upon reaching the desired temperature, the crucible was held at this temperature for 5 min. Then the crucible was quenched directly into cold water to rapidly solidify the molten metal and prevent further particle precipitation. The quenched sample was sectioned, polished and observed under optical and electronic microscopes to detect the possible presence of intermetallic particles. The set-up of the LiMCA detection is shown in Fig. 1 (Fortier and Chen, 2004). A LiMCA tube with a 300 m orifice was placed under the melt surface for the continuous monitoring of the foreign particle number and size in the liquid aluminum. A thermocouple was inserted into the melt beside the LiMCA tube to monitor the melt temperature during the test. From 770 ◦ C, the melt temperature was gradually lowered at a rate of approximately 0.5 ◦ C/min. The particle number, expressed as N20 count, was measured with time. The N20 count represents the amount of foreign particles with a size larger than 20 m inside the melt in thousands per kg of metal. During the slow decrease of the melt temperature, a sharp rise of the particle count N20 occurred at a certain temperature. This is considered as the precipitation temperature of the titanium aluminides (Fig. 2). For the determination of the effect of the Ti level and cooling rate on the intermetallic particle quantity and morphology, a set of liquid samples with various Ti contents were poured into the small crucible (∼100 g) and were solidified at different cooling rates. Samples were then metallographically prepared to quantify the titanium aluminide particles.
Al–4.5%Si Al–7%Si (A356) Al–9.5%Si
%Si
%Mg
%Ti
%Fe
4.5 7.0 9.5
0.35 0.35 0.35
0.10–0.26 0.10–0.26 0.10–0.26
0.10 0.10 0.10
2. Experimental procedures
Fig. 1. The set-up of the LiMCA detection.
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Fig. 5. TiAiSi precipitation temperature for Al–9.5%Si alloy.
Fig. 2. Typical LiMCA measurement curve of an Al–7%–0.15%Ti alloy.
It can be seen from the experimental data that the solubility limit of Ti lays around 0.10% in all three Al–Si alloys. It is almost independent on the Si contents in the range from 4.5% to 9.5%. It is evident that the Ti solubility in ternary Al–Si–Ti is lower than the binary Al–Ti system, in which the solubility is situated between 0.12% and 0.15% (Mondolfo, 1976; McCartney, 1989). In the case that the titanium level in the alloy exceeds the solubility limit, the primary TiAlSi intermetallics form directly from the liquid metal. It is clearly seen that as the Ti level increases, primary TiAlSi intermetallic particles precipitate well above the liquidus temperature of the alloy. It becomes obvious that the titanium addition level plays an important role in the presence of such intermetallic particles. 3.2. The growth kinetic
Fig. 3. TiAiSi precipitation temperature for Al–4.5%Si alloy.
3. Lab results and discussion 3.1. The formation of TiAlSi intermetallics The formation temperature curves of titanium aluminides, experimentally obtained for three alloys, are presented in Figs. 3–5 as a function of titanium content. Most results are validated using both melt quenching and LiMCA measurements. The precipitation temperatures of the titanium aluminide are found to be influenced by the titanium and silicon content of the alloy. As can be seen, as the titanium content of the alloy increases, so does the precipitation temperature. Also, as the silicon content is increased, the slope of the precipitation temperature curve is also increased.
Fig. 4. TiAiSi precipitation temperature for Al–7%Si alloy.
From the LiMCA measurement curves, the growth kinetic of intermetallic particles as a function of the temperature has been analyzed. The sudden rise in the particle count in the curve (Fig. 2) indicates quick nucleation of numerous particles. The nucleation process seems to stop rapidly after approximately 1–2 ◦ C of cooling. The growth and settling of existing particles then becomes dominant. By following the particle size distribution with time after the nucleation, it is possible to follow the growth of these particles. Typical results for an A356 alloy are given in Fig. 6. It is shown that the growth of particle corresponds well with the decreasing melt temperature. A linear relationship between the particle size and the melt temperature was found. The maximum particle size recorded in the LiMCA curves was 70–75 m. This is consistent with an equivalent diameter of 85 m of particles measured metallographically in the solidified LiMCA tubes.
Fig. 6. TiAiSi particle size in function of melt temperature for Al–7%Si–0.15%Ti alloy.
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Fig. 7. Effect of Ti level on TiAlSi particle size.
3.3. The effect of Ti level and cooling rate The effect of the titanium level and the cooling rate on the quantity of TiAlSi particles has been evaluated for the A356 alloy. As can be seen in Figs. 7 and 8, the total amount of TiAlSi particles increases with increasing titanium content and with decreasing cooling rates. It can also be seen that the formation of titanium aluminide particles can be suppressed at lower titanium levels using a high cooling rate. Because of the settling of the TiAlSi particles during holding and cooling, the distribution of the particles in the solidified sample is different in the top, middle and bottom sections (Fig. 8). More intermetallic particles are found towards the bottom of the sample. 3.4. The phase morphology and type In all the examined samples of Al–7%Si–Mg (A356) alloy, most TiAlSi intermetallic particles had either a flake-like or a blocky morphology. On a rare occasion, a few petal-like particles were found in the cast microstructure. It is therefore reasonable to believe that the morphology of TiAlSi particles in ternary Al–Si alloys is similar to that of TiAl3 in binary alloy, which was reported as having three distinct morphologies (Arnberg et al., 1982). Although the particles have the same morphology, the backscattered imaging under SEM revealed that they may have completely different compositions. As shown in Fig. 9 the brighter parts in the picture have high Si (∼37 wt.%) and low Al (∼10 wt.%)
Fig. 9. Back-scattered image of SEM showing flake-like TiAlSi particles with different compositions.
while the darker parts have low Si (∼8 wt.%) and high Al (∼40 wt.%) contents. This fact may suggest that few types of ternary TiAlSi intermetallics co-exist in the as-cast structure. Using the EDS technique in SEM, the composition of a large number of particles in A356 alloy was quantitatively analyzed. Three types of TiAlSi intermetallic particles were detected and grouped with following average compositions. I. TiAl3 type, Ti 36–37% – Al 52–54% – Si 9–11% II. Low Si type, Ti 50–55% – Al 37–43% – Si 7–9% III. High Si Type, Ti 50–55% – Al 5–15% – Si 35–39% It is obvious that the formation of TiAlSi intermetallics in the ternary Al–Si–Ti system is quite complex. Because Al and Si can replace each other over a wide range in all three phases (Perrot, 1990), it results in a large variety of chemical compositions of TiAlSi phases. According to EDS results, it is reasonable to believe that the first type of intermetallic particle is Ti(AlSi)3 intermetallic phase. However, the exact identification of other types of intermetallics has to be verified with other advanced analysis techniques. For this reason, the titanium aluminides in the text are generally referred to as TiAlSi intermetallic particles. 4. Industrial results and discussion 4.1. Industrial applications
Fig. 8. Effect of Ti level on TiAlSi particle density on different sections of the solidified sample.
In recent years, horizontal continuous DC casting has been used to produce a variety of aluminum products such as busbars, T-ingots and extrusion billets. It is found that the Ti level and TiAlSi intermetallic precipitation have an important impact on the occurrence of the feeding blockage and wormhole defect during horizontal DC casting. The causes of both cast process related problems were carefully analyzed. A schematic representation of the horizontal continuous DC casting is shown in Fig. 10. The molten aluminum comes from the furnace via a trough to the tundish. Several tubes (usually in cylindrical form) are installed at the bottom of the tundish and connected to the mold. The liquid metal is fed through the feeding tubes into the mold. The liquid metal is first cooled by contact with the mold and then by water sprays out of the mold. Under steady-state, the molten metal is continuously flowed into the mold and the solidifying ingot moves forward horizontally. The horizontal DC machine can cast non-stop up to several days.
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Fig. 10. Schematic set-up of horizontal continuous casting.
4.2. The feeding blockage The feeding blockage, sometimes referred to as “nozzle blockage” in the continuous casting, occurs around and inside of the feeding tubes of the horizontal continuous DC cast unit (Fig. 10). The feeding blockage is recognized today as one of the major hurdles to achieve long cast duration for the horizontal continuous DC casting of Al–Si alloys. The blockage causes defects of the ingots, a short cast duration and even an unexpected interruption of the cast. There are several mechanisms and causes involved in the feeding blockage in the aluminum horizontal continuous casting (Chen, 2006). One of the most encountered cases in the operation is the precipitation of primary TiAlSi intermetallic particles in the feeding tubes. As Ti addition exceeds its solubility, primary TiAlSi intermetallic particles precipitate in the process at temperatures well above the “liquidus”. At typical casting temperature used in the industry, the intermetallic particles already precipitated and suspended in the liquid may flow through the tundish to the feeding tubes. Some of those intermetallic particles can directly deposit on the feed-
ing tube walls and aggregate together to block the aluminum flow. With decreasing melt temperature from the tundish via the feeding tubes to the ingot, the intermetallics continue to precipitate out from the liquid. They grow directly on the entrapped particles to form a network, accelerating the feeding blockage. Fig. 11 shows blocked feeding tubes after an unexpected interruption of the cast. The cast material was an Al–7%Si (A356) alloy and contained 0.15% Ti, exceeding the Ti solubility. The fracture surface of the spouts was metallic and with many shiny crystals. The microstructure of the clogged material revealed that the tubes were full of TiAlSi intermetallic particles. A majority of aggregated and clustered blocky particles were found near the tube wall (Fig. 11c). Long flake-like particles, which reached up to several centimeters in length, are found towards the tube centre (Fig. 11b). The flakelike particles grew towards the liquid centre and were prone to form a network, leading to the total obstruction of the metal flow. The more Ti there was in the alloy, the larger the amount of intermetallic particles was present. A simple phase calculation in Al–7%Si–Ti alloy illustrates the severity of the problem (see Fig. 4). One ton of liquid metal produces ∼0.3 kg intermetallic particles at temperature near the liquidus for every 0.01% Ti that exceeds the solubility. A 1-day long cast of such alloy at ∼300 t/day generates ∼90 kg intermetallics, which have to pass through the feeding tubes. The industrial experiences confirmed that the higher Ti the Al–Si alloys had, the more frequent the feeding blockage and the shorter the cast duration was. In several cases where the A356 alloys contained 0.15–0.18% Ti, the cast operation interrupted within few hours due to a total blockage of the feeding tubes. When Ti levels were decreased to 0.1% and below, this type of the feeding blockage disappeared completely. 4.3. The wormhole defect An undesirable prolonged porosity can occur occasionally in T-ingots of Al–Si alloys (Fig. 12). This defect is often called wormholes and it can reach up to several meters long. Recent study of Chen (2008) showed that the build up of icicles on the insulat-
Fig. 11. An example of feeding blockage by TiAlSi intermetallic particles, (a) appearance of the blocked spouts in the tundish side, (b) macro section of a clogged tube, (c) microstructure near the tube wall, and (d) microstructure towards the tube centre.
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Fig. 12. A typical wormhole in the horizontal cast ingot: the appearance in the cross-section of T-ingot (a) and a macrograph of the wormhole in the longitudinal section (b).
ing plate is a major cause for wormhole formation (see Fig. 10). During casting, an oil-based lubricant is constantly injected into the corner between the mold and the insulating plate. The oil and moisture (contained in the oil and insulating plate) vaporize at the insulating plate and liquid metal interface. When icicles build up on the insulating plate and grow into the mushy zone of the ingot, the gas vapor can penetrate the mushy zone through the icicle capillary channels and generate gas bubbles at the tip of icicles. As the ingot moves forward, the bubbles elongate progressively in the cast direction, and grow continuously up to several meters long. The icicle build up can be generated by a few sources (Chen, 2008). One of the most frequent sources of icicles in the horizontal casting of Al–Si alloys is the precipitation of TiAlSi intermetallics. Similar to the feeding blockage, it was found that, when the Ti level in A356 alloy exceeds its solubility, the TiAlSi intermetallics
precipitate on the insulating plate. Fig. 13a shows an example of large TiAlSi intermetallics agglomeration on the insulating plate which were revealed after a sudden cast abortion where the liquid ran out the mold after a 2 days cast. This was for an alloy containing 0.13%Ti. Two distinct forms of intermetallic particles were found in the microstructure of the TiAlSi agglomeration. Blocky particles were mainly located close to the insulating plate wall where the intermetallic particles initially precipitated (Fig. 13b). In the direction towards the liquid sump, the dominating particles were large and flake-like (Fig. 13c). They were found to extend approximately 30 mm towards the liquid. They aggregated together to form a network and grew continuously towards liquid and often reached up to several centimeters in length. After the Ti content was reduced to 0.10–0.11% or below, no more TiAlSi intermetallic particles was found on the insulting plates and the wormholes in the cast ingots also disappeared.
Fig. 13. A large TiAlSi intermetallic agglomeration on the insulating plate, (a) macrograph of TiAlSi agglomeration, (b) blocky TiAlSi particles (location A), and (c) flake-like TiAlSi particles (location B).
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product problems both in the primary aluminum cast houses and the aluminum foundries. It is therefore recommended that in most of the Al–Si cast alloys, particularly in the widely used 356/357 alloys, the Ti levels should be below or near to 0.10%. 6. Conclusions
Fig. 14. Typical microstructure of an A356 cast wheel part with 0.15%Ti, showing TiAlSi intermetallic particles in the cast structure.
1. The liquid solubility limit of Ti is around 0.10% in Al–Si alloy, in which the Si ranges from 4.5 to 9.5%. 2. When Ti addition exceeds its solubility limit, the primary TiAlSi intermetallics precipitate at temperatures well above the liquidus temperatures of Al–Si cast alloys. 3. The formation temperature of primary TiAlSi intermetallics increases with increasing titanium content. 4. Few types of primary TiAlSi intermetallics can co-exist in Al–Si cast alloys. Their common morphologies in the as-cast structure are flake-like and blocky. 5. Precipitation of primary TiAlSi intermetallics can cause a series of process and product related problems, such as feeding blockage, wormhole defect and low quality of final shape cast parts. 6. It is recommended that the Ti levels in most of the Al–Si cast alloys, particularly in the widely used 356/357 alloys, should be controlled below or near to 0.10%.
4.4. The quality of the cast parts
Acknowledgements
As mentioned earlier, when the Ti addition in Al–Si alloys goes beyond its solubility, TiAlSi intermetallic particles will precipitate in the casting liquid process lines or directly inside the cast parts. During the pouring of liquid into cast molds, some of those particles, precipitated and already suspended in liquid, may go through the gate system and lie in the cast parts. Moreover, the TiSiAl intermetallics can continue to precipitate inside of the cast part during solidification. Fig. 14 shows insoluble TiSiAl intermetallic particles found in low pressure die cast wheel parts of an A356 alloy containing 0.15% Ti. It was found that, when Ti content was larger than 0.11–0.12%, the size and quantity of intermetallic particles in the cast wheel parts increased rapidly with increasing Ti levels. It is evident that large plate-like TiAlSi intermetallics deteriorates the ductility of the cast part. When located near the surface of the casting, they may cause different surface defects, such as hard spots for machining and initial points for pitting corrosion.
The authors would like to thank the National Sciences and Engineering Research Council of Canada (NSERC) and Rio Tinto Alcan for the financial support. The authors gratefully acknowledge the contribution and support of the colleagues at Rio Tinto Alcan who were involved in the project.
5. Perspective of Ti additions in Al–Si cast alloys Because of the grain refinement effect, Ti was introduced into Al–Si cast alloys long time ago. For historical reason, the Ti level in cast alloys has been kept high, far above its solubility, and commonly in the range of 0.13–0.18%. It is well recognized that the Ti alone as a grain refiner is not efficient in aluminum cast alloys. With the introduction of modern grain refinement techniques and efficient master alloys, the required Ti level to achieve the best grain refining is far below its liquid solubility in the presence of TiB2 . For instance, an excellent grain refinement can be achieved nowadays by an addition of 10–20 ppm of B in the form of Al5Ti1B master alloy. This results in actual Ti concentration in the melt of about one-tenth of the liquid solubility of Ti. In addition, it is confirmed that Ti in aluminum matrix show very little solution hardening and there is no precipitation hardening effect during heat treatment (Smith et al., 2004; Sigworth et al., 2007). Therefore, high Ti levels bring no advantage in term of mechanical properties of final cast parts. It is obvious that an excess of Ti addition causes not only increased alloying cost but also introduces a series of process and
References Arnberg, L., Backerud, L., Klang, H., 1982. Grain refinement of aluminium Part2: intermetallic particles in Al–Ti–B-type master alloys for grain refinement of aluminium. Metal Technology (January), 7–13. Chen, X.-G., 2006. Feeding blockage of aluminum horizontal continuous casting. Mater. Sci. Forum 519–521, 1827–1834. Chen, X.-G., 2008. An uncommon cast defect – wormholes – in aluminum DC casting. In: Hirsch, J., Skrotzki, B., Gottstein, G. (Eds.), Aluminum Alloys—Their Physical and Mechanical Properties, ICAA11. Wiley-VCH Verlag, pp. 484–490. Chen, X.-G., Fortier, M., 2004. Formation of primary TiAlSi intermetallic compounds in Al–Si foundry alloys. In: Nie, J.F., Morton, A.J., Muddle, B.C. (Eds.), Proceedings of the 9th International Conference on Aluminum Alloys, pp. 659–665. Fortier, M., Chen, X.-G., 2004. Monitoring of TiAlSi particle formation and growth in Al–Si alloys using LiMCA. In: Tabereaux, A.T. (Ed.), Light Metals. TMS, pp. 845–849. Guzowski, M.M., Sigworth, G.K., Sentner, D.A., 1987. The role of boron in the grain refinement of aluminium with titanium. Met. Trans. A 18, 603–619. Habibi, H.R., Jafarnia, A., Chen, X.-G., Ghomashchi, R., 2003. AlTiSi intermetallics in A356 Al–Si foundry alloy. In: Masounave, J., Dufour, G. (Eds.), Light Metals, COM, pp. 163–170. McCartney, D.G., 1989. Grain refining of aluminium and its alloys using inoculants. Int. Mater. Rev. 34, 247–260. Mondolfo, L.F., 1976. Aluminium Alloys: Structure and Properties. Butterworths, London, pp. 385–387. Perrot, P., 1990. Aluminium–silicon–titanium. In: Petzow, G., Effenberg, G. (Eds.), Ternary Alloys—A Comprehensive Compendium of Evaluated Constitutional Data and Phase Diagrams, vol. 8. VCH Publishers, New York, pp. 283–290. Sigworth, G.K., Smith, C.L., Easton, M.A., Barresi, J., Kuhn, T.A., 2007. The grain refinement of Al–Si casting alloys. In: Sørlie, M. (Ed.), Light Metals. TMS, pp. 691–696. Smith, C.L., Easton, M.A., Nie, J.F., Zhang, X., Couper, M.J., 2004. The effect of Ti content on mechanical properties of an Al7Si0.35Mg alloy. In: Nie, J.F., Morton, A.J., Muddle, B.C. (Eds.), Proceedings of the 9th International Conference on Aluminum Alloys, pp. 1222–1228. Sokolowski, J.H., Kierkus, C.A., Brosnan, B., Evans, W., 2000. Formation of insoluble Ti(AlSi)3 crystals in 356 alloy castings and their sedimentation in foundry equipment: causes, effects and solutions. AFS Trans. 108, 491–495. Spittle, J.A., 2006. Grain refinement in shape casting of aluminium alloys. Int. J. Cast Metals Res. 19, 210–222.
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Journal of Materials Processing Technology journal homepage: www.elsevier.com/locate/jmatprotec
TiAlSi intermetallic formation and its impact on the casting processing in Al–Si alloys X.-G. Chen a,∗ , M. Fortier b a b
University of Québec at Chicoutimi, Chicoutimi, QC, Canada G7H 2B1 Arvida Research and Development Centre, Rio Tinto Alcan Inc., Jonquière, QC, Canada G7S 4K8
a r t i c l e
i n f o
Article history: Received 29 April 2010 Received in revised form 10 June 2010 Accepted 11 June 2010
Keywords: TiAlSi intermetallics Al–Si alloys Microstructure Grain refinement Feeding blockage Wormhole defect
a b s t r a c t The effect of titanium additions on the formation of primary TiAlSi intermetallics and resulted cast microstructure in Al–Si alloys has been studied. The precipitation temperature, growth kinetic and morphology of TiAlSi intermetallics as well as the solubility limit of Ti in Al–Si alloys were experimentally evaluated. The impact of primary TiAlSi intermetallic particles on the casting process related problems was investigated on the large scale of industrial production. The influence of Ti in the occurrence of feeding blockage and wormhole defect during ingot casting was described. The perspectives of Ti addition in Al–Si cast alloys were discussed. It is proposed that the Ti level in the most Al–Si cast alloys should be controlled within 0.10 wt.%. © 2010 Elsevier B.V. All rights reserved.
1. Introduction
aluminum, 665 ◦ C (Mondolfo, 1976):
Aluminum–silicon cast alloys are widely used for the shape casting because of their good castability and mechanical properties. It is a common practice in many aluminum foundries to add titanium in Al–Si cast alloys as a grain refining agent. Although this practice has been employed for many years, Ti addition and grain refining are still a matter of concern to both ingot suppliers and shape casting foundries because of the lack of a definitive knowledge of what is a necessary amount of Ti level to produce adequate grain refining and metallurgical properties. The Ti addition and grain refining practice vary considerable from foundry to foundry, possible ranging from no Ti addition at all to a systematic Ti addition in all cast batches to maintain a uniform in-house operation procedure.
L + TiAl3 ⇔ ˛-Al
1.1. Binary Al–Ti system Before considering the role of Ti in Al–Si alloys, it is useful to first review the binary Al–Ti system. On the aluminum-rich part, Al–Ti alloy presents a peritectic reaction at a composition of approximately 1.2% Ti1 and a temperature above the melting point of pure
∗ Corresponding author at: Department of Applied Sciences, University of Québec at Chicoutimi, Chicoutimi, QC, Canada G7H 2B1. Tel.: +1 418 545 5011x2603; fax: +1 418 545 5012. E-mail address: xgrant [email protected] (X.-G. Chen). 1 Alloy and phase compositions are in wt.% unless specified otherwise. 0924-0136/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2010.06.009
This reaction has a profound influence on nucleation and growth of ␣-aluminum crystals from the liquid stage. TiAl3 intermetallics have an existence range of 36.5–37.5% Ti and a 3370 kg/m3 density (Mondolfo, 1976). The solubility limit of titanium in liquid aluminum is situated between 0.12% and 0.15% Ti (Mondolfo, 1976; McCartney, 1989). Titanium is usually added in the aluminum melt in the form of Al–(5–10%)Ti master alloys. The grain refining is produced by release of numerous TiAl3 intermetallic particles from the master alloys. It is known that these TiAl3 intermetallic particles can act as nucleants for ␣-aluminum crystals during solidification (Arnberg et al., 1982). The formation of TiAl3 intermetallics were studied in the binary Al–Ti system and it was found that TiAl3 intermetallics could have three different morphologies (flakes, petals and blocks) depending on the solidification conditions and the temperature history of the alloy (Arnberg et al., 1982; Guzowski et al., 1987).
1.2. Ternary Al–Si–Ti system At the aluminum-rich corner of ternary Al–Si–Ti system, three possible types of titanium aluminides could be present (Perrot, 1990).
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1. TiAl3 . Up to 15 at.% Al can be replaced by Si in TiAl3 lattice structures, resulting in various chemical compositions and a range of lattice parameters (Mondolfo, 1976). It can be commonly written as Ti(AlSi)3 . 2. 1 : commonly written as Ti7 Al5 Si12 . This phase is stable below 900 ◦ C and can contain 17 up to 42%Si. 3. 2 : commonly written as Ti(AlSi)2 . This phase has higher amount of silicon, which ranges from 38% to 46%.
Table 1 Chemical composition of the investigated alloys (wt.%).
In the most useable Ti ranges of 0–0.3%, all three ternary phases are very close to each other (Perrot, 1990). It can be expected that two or three intermetallic phases can co-exist with aluminum in Al–Si alloys. Because Al and Si can replace each other, it results in a large variety of composition combination and a wide range of lattice parameters of TiAlSi phases. Considering that very little information exists in the literature and equilibriums are hardly achieved in the real casting condition, these titanium aluminides in Al–Si alloys are generally referred to here as TiAlSi intermetallic phases. Habibi et al. (2003) and Fortier and Chen (2004) have recently reported that the morphology of TiAlSi intermetallics in Al–Si foundry alloys is similar to TiAl3 in binary alloy, i.e. flakes and blocks. However, the favourite conditions under which these morphologies would form are still unknown. When Ti concentration added in the melt is less than the solubility limit, TiAlSi intermetallic particles dissolve with time and the grain refining fades away. To avoid this effect, a large amount of Ti (exceeding the solubility) was added for longer lasting refining effect. It is believed that this traditional practice built the basis of Aluminum Association’s chemical composition limits in Al–Si cast alloys many years ago (Sigworth et al., 2007). Although a minimum Ti concentration is not specified, the maximum Ti levels allowed in most Al–Si alloys are very high, mostly in the range of 0.2–0.25%. A number of studies have shown that Ti alone is a relatively weak grain refiner in aluminum alloys in despite of maintaining high Ti level in the melt (Arnberg et al., 1982; Guzowski et al., 1987). The addition of boron and titanium in the form of Al–Ti–B master alloys has been found to be much more effective in grain refinement than titanium alone, since boron in the form of TiB2 enhances significantly the grain refining action. Different Al–Ti–B master alloys with various Ti/B ratios, such as Al5Ti1B and Al3Ti1B, have been developed to improve significantly the grain refinement efficiency in Al–Si cast alloys (Spittle, 2006; Sigworth et al., 2007). An excess of Ti addition can cause a number of problems during liquid metal processing and defects in the final cast products. Sigworth et al. (2007) reported that when a high Ti melt is held in the furnace, TiAlSi crystals precipitated and formed a layer of “sludge” at the bottom of the furnace. Studies of Sokolowski et al. (2000) and Chen and Fortier (2004) revealed that a high level of Ti in the melt created massive TiAlSi crystal precipitation which caused a heavy sedimentation in liquid metal processing lines and a presence of coarse intermetallic crystals in the cast structure. These crystals proved to be detrimental to the foundry processes and the metallurgical properties of the castings. Finally, aluminum foundries obtain their base Al–Si alloys from different sources. The use of pre-inoculated (with Ti) cast ingots is still one important source. In addition, a significant proportion of foundry alloys are in-house recycled materials or mixed with secondary alloys, which usually contained a certain amount of residual Ti. The addition of an excessive Ti combined with non-controlled residual Ti can result in the precipitation of primary TiAlSi intermetallic particles. Therefore, it is of great interest to establish the optimum titanium levels for Al–Si alloys, which can achieve the most desirable metallurgical properties. The aim of the present study is to examine the impact of Ti and its TiAlSi intermetallics on the casting processing in Al–Si cast alloys.
A batch of 5–20 kg of commercial purity aluminum was melted in an electric resistance furnace and alloyed to the desired compositions at 750–770 ◦ C. Three Al–Si cast alloys were studied at different titanium levels ranging from 0.10% to 0.26% and their chemical compositions are given in Table 1. Titanium additions in the melt were achieved using a binary Al–6%Ti master alloy. The evaluation of the formation temperature of titanium aluminides was done using the melt quenching and the LiMCA (Liquid Metal Cleanliness Analyser) detection. In the melt quenching tests, a certain amount of the molten metal (∼100 g) was poured into a small crucible and placed into a furnace at 800 ◦ C. The furnace temperature was gradually lowered to reach the desired temperature. Upon reaching the desired temperature, the crucible was held at this temperature for 5 min. Then the crucible was quenched directly into cold water to rapidly solidify the molten metal and prevent further particle precipitation. The quenched sample was sectioned, polished and observed under optical and electronic microscopes to detect the possible presence of intermetallic particles. The set-up of the LiMCA detection is shown in Fig. 1 (Fortier and Chen, 2004). A LiMCA tube with a 300 m orifice was placed under the melt surface for the continuous monitoring of the foreign particle number and size in the liquid aluminum. A thermocouple was inserted into the melt beside the LiMCA tube to monitor the melt temperature during the test. From 770 ◦ C, the melt temperature was gradually lowered at a rate of approximately 0.5 ◦ C/min. The particle number, expressed as N20 count, was measured with time. The N20 count represents the amount of foreign particles with a size larger than 20 m inside the melt in thousands per kg of metal. During the slow decrease of the melt temperature, a sharp rise of the particle count N20 occurred at a certain temperature. This is considered as the precipitation temperature of the titanium aluminides (Fig. 2). For the determination of the effect of the Ti level and cooling rate on the intermetallic particle quantity and morphology, a set of liquid samples with various Ti contents were poured into the small crucible (∼100 g) and were solidified at different cooling rates. Samples were then metallographically prepared to quantify the titanium aluminide particles.
Al–4.5%Si Al–7%Si (A356) Al–9.5%Si
%Si
%Mg
%Ti
%Fe
4.5 7.0 9.5
0.35 0.35 0.35
0.10–0.26 0.10–0.26 0.10–0.26
0.10 0.10 0.10
2. Experimental procedures
Fig. 1. The set-up of the LiMCA detection.
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Fig. 5. TiAiSi precipitation temperature for Al–9.5%Si alloy.
Fig. 2. Typical LiMCA measurement curve of an Al–7%–0.15%Ti alloy.
It can be seen from the experimental data that the solubility limit of Ti lays around 0.10% in all three Al–Si alloys. It is almost independent on the Si contents in the range from 4.5% to 9.5%. It is evident that the Ti solubility in ternary Al–Si–Ti is lower than the binary Al–Ti system, in which the solubility is situated between 0.12% and 0.15% (Mondolfo, 1976; McCartney, 1989). In the case that the titanium level in the alloy exceeds the solubility limit, the primary TiAlSi intermetallics form directly from the liquid metal. It is clearly seen that as the Ti level increases, primary TiAlSi intermetallic particles precipitate well above the liquidus temperature of the alloy. It becomes obvious that the titanium addition level plays an important role in the presence of such intermetallic particles. 3.2. The growth kinetic
Fig. 3. TiAiSi precipitation temperature for Al–4.5%Si alloy.
3. Lab results and discussion 3.1. The formation of TiAlSi intermetallics The formation temperature curves of titanium aluminides, experimentally obtained for three alloys, are presented in Figs. 3–5 as a function of titanium content. Most results are validated using both melt quenching and LiMCA measurements. The precipitation temperatures of the titanium aluminide are found to be influenced by the titanium and silicon content of the alloy. As can be seen, as the titanium content of the alloy increases, so does the precipitation temperature. Also, as the silicon content is increased, the slope of the precipitation temperature curve is also increased.
Fig. 4. TiAiSi precipitation temperature for Al–7%Si alloy.
From the LiMCA measurement curves, the growth kinetic of intermetallic particles as a function of the temperature has been analyzed. The sudden rise in the particle count in the curve (Fig. 2) indicates quick nucleation of numerous particles. The nucleation process seems to stop rapidly after approximately 1–2 ◦ C of cooling. The growth and settling of existing particles then becomes dominant. By following the particle size distribution with time after the nucleation, it is possible to follow the growth of these particles. Typical results for an A356 alloy are given in Fig. 6. It is shown that the growth of particle corresponds well with the decreasing melt temperature. A linear relationship between the particle size and the melt temperature was found. The maximum particle size recorded in the LiMCA curves was 70–75 m. This is consistent with an equivalent diameter of 85 m of particles measured metallographically in the solidified LiMCA tubes.
Fig. 6. TiAiSi particle size in function of melt temperature for Al–7%Si–0.15%Ti alloy.
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Fig. 7. Effect of Ti level on TiAlSi particle size.
3.3. The effect of Ti level and cooling rate The effect of the titanium level and the cooling rate on the quantity of TiAlSi particles has been evaluated for the A356 alloy. As can be seen in Figs. 7 and 8, the total amount of TiAlSi particles increases with increasing titanium content and with decreasing cooling rates. It can also be seen that the formation of titanium aluminide particles can be suppressed at lower titanium levels using a high cooling rate. Because of the settling of the TiAlSi particles during holding and cooling, the distribution of the particles in the solidified sample is different in the top, middle and bottom sections (Fig. 8). More intermetallic particles are found towards the bottom of the sample. 3.4. The phase morphology and type In all the examined samples of Al–7%Si–Mg (A356) alloy, most TiAlSi intermetallic particles had either a flake-like or a blocky morphology. On a rare occasion, a few petal-like particles were found in the cast microstructure. It is therefore reasonable to believe that the morphology of TiAlSi particles in ternary Al–Si alloys is similar to that of TiAl3 in binary alloy, which was reported as having three distinct morphologies (Arnberg et al., 1982). Although the particles have the same morphology, the backscattered imaging under SEM revealed that they may have completely different compositions. As shown in Fig. 9 the brighter parts in the picture have high Si (∼37 wt.%) and low Al (∼10 wt.%)
Fig. 9. Back-scattered image of SEM showing flake-like TiAlSi particles with different compositions.
while the darker parts have low Si (∼8 wt.%) and high Al (∼40 wt.%) contents. This fact may suggest that few types of ternary TiAlSi intermetallics co-exist in the as-cast structure. Using the EDS technique in SEM, the composition of a large number of particles in A356 alloy was quantitatively analyzed. Three types of TiAlSi intermetallic particles were detected and grouped with following average compositions. I. TiAl3 type, Ti 36–37% – Al 52–54% – Si 9–11% II. Low Si type, Ti 50–55% – Al 37–43% – Si 7–9% III. High Si Type, Ti 50–55% – Al 5–15% – Si 35–39% It is obvious that the formation of TiAlSi intermetallics in the ternary Al–Si–Ti system is quite complex. Because Al and Si can replace each other over a wide range in all three phases (Perrot, 1990), it results in a large variety of chemical compositions of TiAlSi phases. According to EDS results, it is reasonable to believe that the first type of intermetallic particle is Ti(AlSi)3 intermetallic phase. However, the exact identification of other types of intermetallics has to be verified with other advanced analysis techniques. For this reason, the titanium aluminides in the text are generally referred to as TiAlSi intermetallic particles. 4. Industrial results and discussion 4.1. Industrial applications
Fig. 8. Effect of Ti level on TiAlSi particle density on different sections of the solidified sample.
In recent years, horizontal continuous DC casting has been used to produce a variety of aluminum products such as busbars, T-ingots and extrusion billets. It is found that the Ti level and TiAlSi intermetallic precipitation have an important impact on the occurrence of the feeding blockage and wormhole defect during horizontal DC casting. The causes of both cast process related problems were carefully analyzed. A schematic representation of the horizontal continuous DC casting is shown in Fig. 10. The molten aluminum comes from the furnace via a trough to the tundish. Several tubes (usually in cylindrical form) are installed at the bottom of the tundish and connected to the mold. The liquid metal is fed through the feeding tubes into the mold. The liquid metal is first cooled by contact with the mold and then by water sprays out of the mold. Under steady-state, the molten metal is continuously flowed into the mold and the solidifying ingot moves forward horizontally. The horizontal DC machine can cast non-stop up to several days.
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Fig. 10. Schematic set-up of horizontal continuous casting.
4.2. The feeding blockage The feeding blockage, sometimes referred to as “nozzle blockage” in the continuous casting, occurs around and inside of the feeding tubes of the horizontal continuous DC cast unit (Fig. 10). The feeding blockage is recognized today as one of the major hurdles to achieve long cast duration for the horizontal continuous DC casting of Al–Si alloys. The blockage causes defects of the ingots, a short cast duration and even an unexpected interruption of the cast. There are several mechanisms and causes involved in the feeding blockage in the aluminum horizontal continuous casting (Chen, 2006). One of the most encountered cases in the operation is the precipitation of primary TiAlSi intermetallic particles in the feeding tubes. As Ti addition exceeds its solubility, primary TiAlSi intermetallic particles precipitate in the process at temperatures well above the “liquidus”. At typical casting temperature used in the industry, the intermetallic particles already precipitated and suspended in the liquid may flow through the tundish to the feeding tubes. Some of those intermetallic particles can directly deposit on the feed-
ing tube walls and aggregate together to block the aluminum flow. With decreasing melt temperature from the tundish via the feeding tubes to the ingot, the intermetallics continue to precipitate out from the liquid. They grow directly on the entrapped particles to form a network, accelerating the feeding blockage. Fig. 11 shows blocked feeding tubes after an unexpected interruption of the cast. The cast material was an Al–7%Si (A356) alloy and contained 0.15% Ti, exceeding the Ti solubility. The fracture surface of the spouts was metallic and with many shiny crystals. The microstructure of the clogged material revealed that the tubes were full of TiAlSi intermetallic particles. A majority of aggregated and clustered blocky particles were found near the tube wall (Fig. 11c). Long flake-like particles, which reached up to several centimeters in length, are found towards the tube centre (Fig. 11b). The flakelike particles grew towards the liquid centre and were prone to form a network, leading to the total obstruction of the metal flow. The more Ti there was in the alloy, the larger the amount of intermetallic particles was present. A simple phase calculation in Al–7%Si–Ti alloy illustrates the severity of the problem (see Fig. 4). One ton of liquid metal produces ∼0.3 kg intermetallic particles at temperature near the liquidus for every 0.01% Ti that exceeds the solubility. A 1-day long cast of such alloy at ∼300 t/day generates ∼90 kg intermetallics, which have to pass through the feeding tubes. The industrial experiences confirmed that the higher Ti the Al–Si alloys had, the more frequent the feeding blockage and the shorter the cast duration was. In several cases where the A356 alloys contained 0.15–0.18% Ti, the cast operation interrupted within few hours due to a total blockage of the feeding tubes. When Ti levels were decreased to 0.1% and below, this type of the feeding blockage disappeared completely. 4.3. The wormhole defect An undesirable prolonged porosity can occur occasionally in T-ingots of Al–Si alloys (Fig. 12). This defect is often called wormholes and it can reach up to several meters long. Recent study of Chen (2008) showed that the build up of icicles on the insulat-
Fig. 11. An example of feeding blockage by TiAlSi intermetallic particles, (a) appearance of the blocked spouts in the tundish side, (b) macro section of a clogged tube, (c) microstructure near the tube wall, and (d) microstructure towards the tube centre.
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Fig. 12. A typical wormhole in the horizontal cast ingot: the appearance in the cross-section of T-ingot (a) and a macrograph of the wormhole in the longitudinal section (b).
ing plate is a major cause for wormhole formation (see Fig. 10). During casting, an oil-based lubricant is constantly injected into the corner between the mold and the insulating plate. The oil and moisture (contained in the oil and insulating plate) vaporize at the insulating plate and liquid metal interface. When icicles build up on the insulating plate and grow into the mushy zone of the ingot, the gas vapor can penetrate the mushy zone through the icicle capillary channels and generate gas bubbles at the tip of icicles. As the ingot moves forward, the bubbles elongate progressively in the cast direction, and grow continuously up to several meters long. The icicle build up can be generated by a few sources (Chen, 2008). One of the most frequent sources of icicles in the horizontal casting of Al–Si alloys is the precipitation of TiAlSi intermetallics. Similar to the feeding blockage, it was found that, when the Ti level in A356 alloy exceeds its solubility, the TiAlSi intermetallics
precipitate on the insulating plate. Fig. 13a shows an example of large TiAlSi intermetallics agglomeration on the insulating plate which were revealed after a sudden cast abortion where the liquid ran out the mold after a 2 days cast. This was for an alloy containing 0.13%Ti. Two distinct forms of intermetallic particles were found in the microstructure of the TiAlSi agglomeration. Blocky particles were mainly located close to the insulating plate wall where the intermetallic particles initially precipitated (Fig. 13b). In the direction towards the liquid sump, the dominating particles were large and flake-like (Fig. 13c). They were found to extend approximately 30 mm towards the liquid. They aggregated together to form a network and grew continuously towards liquid and often reached up to several centimeters in length. After the Ti content was reduced to 0.10–0.11% or below, no more TiAlSi intermetallic particles was found on the insulting plates and the wormholes in the cast ingots also disappeared.
Fig. 13. A large TiAlSi intermetallic agglomeration on the insulating plate, (a) macrograph of TiAlSi agglomeration, (b) blocky TiAlSi particles (location A), and (c) flake-like TiAlSi particles (location B).
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product problems both in the primary aluminum cast houses and the aluminum foundries. It is therefore recommended that in most of the Al–Si cast alloys, particularly in the widely used 356/357 alloys, the Ti levels should be below or near to 0.10%. 6. Conclusions
Fig. 14. Typical microstructure of an A356 cast wheel part with 0.15%Ti, showing TiAlSi intermetallic particles in the cast structure.
1. The liquid solubility limit of Ti is around 0.10% in Al–Si alloy, in which the Si ranges from 4.5 to 9.5%. 2. When Ti addition exceeds its solubility limit, the primary TiAlSi intermetallics precipitate at temperatures well above the liquidus temperatures of Al–Si cast alloys. 3. The formation temperature of primary TiAlSi intermetallics increases with increasing titanium content. 4. Few types of primary TiAlSi intermetallics can co-exist in Al–Si cast alloys. Their common morphologies in the as-cast structure are flake-like and blocky. 5. Precipitation of primary TiAlSi intermetallics can cause a series of process and product related problems, such as feeding blockage, wormhole defect and low quality of final shape cast parts. 6. It is recommended that the Ti levels in most of the Al–Si cast alloys, particularly in the widely used 356/357 alloys, should be controlled below or near to 0.10%.
4.4. The quality of the cast parts
Acknowledgements
As mentioned earlier, when the Ti addition in Al–Si alloys goes beyond its solubility, TiAlSi intermetallic particles will precipitate in the casting liquid process lines or directly inside the cast parts. During the pouring of liquid into cast molds, some of those particles, precipitated and already suspended in liquid, may go through the gate system and lie in the cast parts. Moreover, the TiSiAl intermetallics can continue to precipitate inside of the cast part during solidification. Fig. 14 shows insoluble TiSiAl intermetallic particles found in low pressure die cast wheel parts of an A356 alloy containing 0.15% Ti. It was found that, when Ti content was larger than 0.11–0.12%, the size and quantity of intermetallic particles in the cast wheel parts increased rapidly with increasing Ti levels. It is evident that large plate-like TiAlSi intermetallics deteriorates the ductility of the cast part. When located near the surface of the casting, they may cause different surface defects, such as hard spots for machining and initial points for pitting corrosion.
The authors would like to thank the National Sciences and Engineering Research Council of Canada (NSERC) and Rio Tinto Alcan for the financial support. The authors gratefully acknowledge the contribution and support of the colleagues at Rio Tinto Alcan who were involved in the project.
5. Perspective of Ti additions in Al–Si cast alloys Because of the grain refinement effect, Ti was introduced into Al–Si cast alloys long time ago. For historical reason, the Ti level in cast alloys has been kept high, far above its solubility, and commonly in the range of 0.13–0.18%. It is well recognized that the Ti alone as a grain refiner is not efficient in aluminum cast alloys. With the introduction of modern grain refinement techniques and efficient master alloys, the required Ti level to achieve the best grain refining is far below its liquid solubility in the presence of TiB2 . For instance, an excellent grain refinement can be achieved nowadays by an addition of 10–20 ppm of B in the form of Al5Ti1B master alloy. This results in actual Ti concentration in the melt of about one-tenth of the liquid solubility of Ti. In addition, it is confirmed that Ti in aluminum matrix show very little solution hardening and there is no precipitation hardening effect during heat treatment (Smith et al., 2004; Sigworth et al., 2007). Therefore, high Ti levels bring no advantage in term of mechanical properties of final cast parts. It is obvious that an excess of Ti addition causes not only increased alloying cost but also introduces a series of process and
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